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CeramicArmorMaterials by Design
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Ceramic ArmorMaterials by Design
Volume 134
CeramicTransactions
Published byThe American Ceramic Society735 Ceramic PlaceWesterville, Ohio 43081www.ceramics.org
Edited by
James W. McCauleyU.S. Army Research Laboratory
Andrew Crowson U.S. Army Research Laboratory
William A. Gooch, Jr.U.S. Army Research Laboratory
A.M. RajendranU.S. Army Research Laboratory
Stephan J. BlessThe University of Texas at Austin
Kathryn V. LoganGeorgia Institute of Technology
Michael NormandiaU.S. Army Research Laboratory
Steven WaxU.S. Defense Advanced Research Projects Agency
Proceedings of the Ceramic Armor Materials by Design Symposium held at the Pac RimIV International Conference on Advanced Ceramics and Glass, November 4–8, 2001 inWailea, Maui, Hawaii.
Proceedings of the Ceramic Armor Materials by Design Symposium held at the Pac Rim IVInternational Conference on Advanced Ceramics and Glass, November 4–8, 2001 in Wailea, Maui,Hawaii.
Copyright 2002,The American Ceramic Society. All rights reserved.
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COVER PHOTO:“Post-test photograph of impact of AP bullet against ceramic/aluminum tar-get: (a) front view of ceramic element;” is courtesy of Charles E. Anderson Jr., and appears asfigure 4a in his paper “Developing an Ultra-Lightweight Armor Concept,” which begins on page485.
For information on ordering titles published by The American Ceramic Society, or to request apublications catalog, please call 614-794-5890.
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4 3 2 1–05 04 03 02
ISSN 1042-1122 ISBN 1-57498-148-X
v
Preface . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . xi
CERAMIC ARMOR DEVELOPMENTAn Overview of Ceramic Armor Applications . . . . . . . . . . . . . . . 3
William A. Gooch Jr., U.S. Army Research Laboratory
Armor Ceramics Under High-Velocity Impact of a Medium-Caliber Long-Rod Penetrator. . . . . . . . . . . . . . . . . . . . 23
Hans-Jürgen Ernst,Volker Wiesner, and Thomas Wolf,French-German Research Institute of Saint-Louis (ISL)
Practical Issues in Ceramic Armor Design . . . . . . . . . . . . . . . . . 33Bryn James, Defense Science and Technology Laboratories
Ballistic Development of Tungsten Carbide Ceramics for Armor Applications . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 45
Pierre-François Peron, Etablissement Technique de Bourges
Ballistic Development of U.S. High Density Tungsten Carbide Ceramics . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 53
William A. Gooch and Matthew S. Burkins, U.S. Army Research Laboratory; Richard Palicka, Cercom Incorporated
Initial Tests on Ceramics in Composite Armor . . . . . . . . . . . . . 63W. Lanz, RUAG Land Systems
Structure and Properties of Shock-Resistant Ceramics Developed at the Institute for Problems in Materials Science . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 73
B.A. Galanov, O.N. Grigoriev, S.M. Ivanov, and V.V. Kartuzov, National Academy of Sciences of Ukraine
Ceramic Armor with Submicron Alumina Against Armor Piercing Projectiles. . . . . . . . . . . . . . . . . . . . . . . . . . . . . 83
E. Strassburger and B. Lexow, Fraunhofer-Institut für Kurzzeitdynamik Ernst-Mach-Institut (EMI); A. Krell, Fraunhofer-Institut für Keramische Technologien und Sinterwerkstoffe
Contents
vi
Armor Alumina Ceramics . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 91Eugene Medvedovski, Ceramic Protection Corporation
Ballistic Performance of Alumina Ceramic Armors. . . . . . . . . . 103Murat Vural and Zeki Erim, Istanbul Technical University;B.A. Konduk and A.H. Ucisik, Bogazici University
PENETRATION AND BALLISTIC TESTINGAn Overview of Ballistic Testing Methods of Ceramic Materials . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 113
Michael J. Normandia and William A. Gooch, U.S. Army Research Laboratory
Theory and Experimental Test Methods for Evaluating Ceramic Armor Components . . . . . . . . . . . . . . . . . 139
Marc A. Adams, Jet Propulsion Laboratory
Long Rod Penetration of Ceramics . . . . . . . . . . . . . . . . . . . . . 151D.L. Orphal, International Research Associates
Depth of Penetration Testing . . . . . . . . . . . . . . . . . . . . . . . . . . 165Bryn James, Defense Science and Technology Laboratories
Transition Between Interface Defeat and Penetration for a Given Combination of Projectile and Ceramic Material . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 173
Patrik Lundberg, René Renström, and Lars Westerling, Swedish Defense Research Agency, FOI
SHOCK AND HIGH STRAIN RATE DYNAMICDynamic Fracture of Ceramics and CMC . . . . . . . . . . . . . . . . 185
Albert S. Kobayashi, University of Washington
Compressive Fracture of Brittle Solids Under Shock-Wave Loading . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 197
G. I. Kanel, Institute for High Energy Densities; S. J. Bless, The University of Texas at Austin
Recent Developments in Split Hopkinson Pressure Bar Testing. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 217
W. Chen and B. Song, The University of Arizona; D. J. Frew and M. J. Forrestal, Sandia National Laboratories
vii
Using Bar Impact to Determine Dynamic Properties of Ceramics . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 225
Stephan J. Bless, The University of Texas at Austin
Shock Compression and Release Properties of Coors AD995 Alumina. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 233
William D. Reinhart and Lalit C. Chhabildas, Sandia National Laboratories; Dennis E. Grady, Applied Research Associates; andTsutomu Mashimo, Kumamoto University
Compressibility and Shear Strength of Titanium Diboride Under Plane Shock Wave Loading . . . . . . . . . . . . . . . . . . . . . . 249
D. P. Dandekar and E. J. Rapacki, U.S. Army Research Laboratory
Dynamic Indentation Damage of Ceramics . . . . . . . . . . . . . . . 261Do Kyung Kim, Chul-Seung Lee, and Young-Gu Kim, Korea Advanced Institute of Science and Technology; Chang Wook Kim and Soon Nam Chang, Agency for Defense Development
Taylor-Impact Experiments for Brittle Ceramic Materials. . . . . 269L. C. Chhabildas and W. D. Reinhart, Sandia National Laboratories;D. P. Dandekar, U.S. Army Research Laboratory
ANALYTICAL AND COMPUTATIONAL MODELINGHistorical Perspective on Ceramic Materials Damage Models . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 281
A.M. Rajendran, U.S. Army Research Laboratory
A Comparison of Ceramic Material Models . . . . . . . . . . . . . . 299Douglas W.Templeton, U. S. Army Tank Automotive Research,Development, and Engineering Center;Timothy J. Holmquist,Network Computing Services Inc./Army HPC Research Center;Hubert W. Meyer Jr., David J. Grove, and Brian Leavy, U.S. Army Research Laboratory
Modeling Ceramic Dwell and Interface Defeat . . . . . . . . . . . . 309Timothy J. Holmquist and Gordon R. Johnson, Network CS/Army High Performance Computing Research Center
3D Finite Element Analysis of Impact Damage in Metallic and Ceramic Targets . . . . . . . . . . . . . . . . . . . . . . . . . . 317
Fenghua Zhou and Jean-Francois Molinari, Johns Hopkins University
viii
A Numerical Investigation of Microcracking Diffusion in Sandwiched Glass Plates . . . . . . . . . . . . . . . . . . . . 329
Z. Chen and L. Shen, University of Missouri-Columbia; G.I. Kanel and S.V. Razorenov, Russian Academy of Sciences
Analytic Model for Penetration of Thick Ceramic Targets . . . . 337James D.Walker, Southwest Research Institute
Grain Level Analysis of Ceramic Microstructures Subjected to Impact Loading . . . . . . . . . . . . . . . . . . . . . . . . . . 349
Pablo D. Zavattieri and Horacio D. Espinosa, Northwestern University
Analysis and Modeling of Ceramic Armor Penetration. . . . . . . 361S.J. Cimpoeru and R.L.Woodward, DSTO Aeronautical and Maritime Research Laboratory
Overview of the Rajendran-Grove Ceramic Failure Model . . . 371D. J. Grove and A. M. Rajendran, U. S. Army Research Laboratory
DAMAGE EVOLUTION AND MICROMECHANISMSFailure Phenomenology of Confined Ceramic Targets and Impacting Rods . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 385
Donald A. Shockey and A.H. Marchand, SRI International;S.R. Skaggs, G.E. Cort, M.W. Burkett, and R. Parker, Los Alamos National Laboratory
Micro-Mechanisms of Compression Failure . . . . . . . . . . . . . . . 403Sia Nemat-Nasser and Sai Sarva, University of California, San Diego
Damage Mitigation in Ceramics: Historical Developments and Future Directions in Army Research . . . . . . . . . . . . . . . . . 421
D.M. Stepp, U.S. Army Research Office
Indentation Damage Behavior of Armor Ceramics. . . . . . . . . . 429Do Kyung Kim and Chul-Seung Lee, Korea Advanced Institute of Science and Technology; Chang Wook Kim and Soon Nam Chang, Agency for Defense Development
Progress in the 3-D Visualization of Interior Ballistic Damage in Armor Ceramics . . . . . . . . . . . . . . . . . . . . . . . . . . 441
Joseph M.Wells, Nevin L. Rupert, and William H. Green, U.S. Army Research Laboratory
ix
PROCESSING AND MANUFACTURINGAn Assessment of Low Cost Manufacturing Technology for Advanced Structural Ceramics and Its Impact on Ceramic Armor . . . . . . . . . . . . . . . . . . . . . . 451
Richard E.Tressler, The Pennsylvania State University
High-Purity Submicron �-Al2O3 Armor Ceramics Design, Manufacture, and Ballistic Performance . . . . . . . . . . . . 463
Andreas Krell, Fraunhofer Institut für Keramische Technologien und Sinterwerkstoffe(IKTS); Elmar Strassburger, Fraunhofer Institut für Kurzzeitdynamik (EMI)
Solid Freeform Fabrication of Advanced Armor Concepts: Opportunities for Design and Manufacture . . . . . . . 473
R.C. McCuiston, S.C. Danforth, M.J. Matthewson, and D.E. Niesz,Rutgers,The State University of New Jersey
ULTRA-LIGHTWEIGHT AND NOVEL CONCEPTSDeveloping an Ultra-Lightweight Armor Concept . . . . . . . . . . 485
Charles E. Anderson Jr., Southwest Research Institute
Ceramics That Exhibit a Threshold Strength . . . . . . . . . . . . . . 499F. F. Lange, M.P. Rao, K. Hbaieb, and R.M. McMeeking,University of California at Santa Barbara
Novel Ideas in Multi-Functional Ceramic Armor Design . . . . . 511Sia Nemat-Nasser, Sai Sarva, Jon B. Isaacs, and David W. Lischer,University of California, San Diego
A New Family of Reaction Bonded Ceramics for Armor Applications . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 527
M. K. Aghajanian, B. N. Morgan, J. R. Singh, M Cubed Technologies, Inc.;J. Mears and R. A.Wolffe, Simula Safety Systems, Inc.
Flexible Ceramic Coated Fiber Fabrics for Lightweight Protection Systems . . . . . . . . . . . . . . . . . . . . . . . . 541
Konstantin von Niessen and Rainer Gadow, University of Stuttgart
Improved Performance of Alumina Ceramics with Carbon Nanotube Reinforcement . . . . . . . . . . . . . . . . . . . . . . 551
Michael Sennett, Natick Soldier Center; Sekyung Chang,Robert H. Doremus, Richard W. Siegel, Pulickel M. Ajayan, and Linda S. Schadler, Rensselaer Polytechnic Institute
x
Recent Progress on the Influence of Microstructure and Mechanical Properties on Ballistic Performance . . . . . . . . 557
J.C. LaSalvia, U.S. Army Research Laboratory
Transparent ArmorTransparent Armor Materials: Needs and Requirements . . . . . 573
Parimal J. Patel and Gary A. Gilde, U.S. Army Research Laboratory
Microwave Reactive Sintering to Fully Transparent Aluminum Oxynitride (AlON) Ceramics . . . . . . . . . . . . . . . . . 587
Dinesh Agrawal, Jiping Cheng, and Rustum Roy, The Pennsylvania State University
An Investigation of the Transmission Properties and Ballistic Performance of Hot Pressed Spinel. . . . . . . . . . . . . . . 595
Mark C.L. Patterson, Technology Assessment & Transfer Inc.; Don W. Roy,Independent; and Gary Gilde, U.S. Army Research Laboratory
Microstructure and Macrostructure EffectsThe Effect of Microstructure on the Dynamic Behavior of Composite Alumina/Titanium Diboride . . . . . . . . . . . . . . . . 611
Kathryn V. Logan, Georgia Institute of Technology
Phase Equilibrium Studies in Al2O3-TiB2 . . . . . . . . . . . . . . . . . . 623Isabel K. Lloyd, University of Maryland; Kevin J. Doherty and Gary A. Gilde, U.S. Army Research Laboratory
Microstructure Development of Aluminum Oxide/TitaniumDiboride Composites for Penetration Resistance . . . . . . . . . . 629
J.W. Adams, G.A. Gilde, and M. Burkins, U.S. Army Research Laboratory; L. Prokurat Franks, U.S. Army Tank-Automotive and Armaments Command
The Effect of Metal-Ceramic Bonding on Ballistic Impact. . . . . 635Kevin J. Doherty, U.S. Army Research Laboratory
Aspects of Geometry Affecting the Ballistic Performance of Ceramic Targets. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 643
I. M. Pickup, A. K. Barker, R. Chenari, and B. J. James, Defense Science and Technology Laboratories;V. Hohler, K.Weber, and R.Tham,Faunhofer-Institut fur Kurzzeitdynamik (EMI)
Index . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 651
xi
This volume contains the proceedings of the “Ceramic Armor Materials byDesign” symposium held at the Pac Rim IV International Conference onAdvanced Ceramics and Glasses held November 4–8, 2001 in Wailea, Maui,Hawaii.
In 1998, the Army formally approved a new basic research StrategicResearch Objective (SRO)—“Armor Materials by Design”. This actionresulted from a critical assessment of the survivability requirements offuture lightweight weapon systems, as well as the emerging materials andmechanics science and engineering that could be brought to bear on thisproblem. It was concluded that there was a critical need for an integrated,multi-disciplinary basic research program that would result in the capabilityto actually design materials for passive, kinetic energy, armor applications.
Since some high performance structural ceramics have been shown to haveoutstanding armor properties at relatively low weight, the symposium wasorganized to address the ceramic armor aspects of the SRO. Researchersfrom around the world working in private industry, academia, and govern-ment organizations on passive transparent and opaque ceramic armorwere invited to participate in this special program.
It was the goal of the symposium to connect ballistic performance tomacro, micro, and crystallographic mechanisms of damage evolution as wellas static and high strain rate mechanical properties and to assess the cur-rent status of computer codes to model and simulate the ballistic per-formance of these materials against kinetic energy projectiles. Currentstate-of-the-art research and development, as well as some historical con-tent, was incorporated into an integrated program.
Most of the credit for this symposium goes to the organizing committeeconsisting of William A. Gooch Jr. and Michael Normandia, U.S. Army
Preface
Research Laboratory, Andrew Crowson, A. M. Rajendran, and David Stepp,Army Research Office of the Army Research Laboratory, Stephan J. Bless,University of Texas, and Steven Wax, Defense Advanced Research ProjectsAgency.
The symposium was co-sponsored by Steven Wax of the U. S. DefenseAdvanced Research Projects Agency, Andrew Crowson, A. M. Rajendranand David Stepp of the Army Research Office of the Army ResearchLaboratory, and William A. Gooch Jr. and James W. McCauley of the U. S.Army Research Laboratory.This support was critical to the success of thesymposium.
Finally, thanks also go to Ms. Susan J. Burns, Battelle, Research Triangle Park,NC for her tremendous help with assembling this book.
James W. McCauley, Chair, Organizing Committee
Editors
James W. McCauley
Andrew Crowson
William A. Gooch, Jr.
M. Rajendran
Stephan J. Bless
Kathryn V. Logan
Michael Normandia
Steven Wax
Ceramic Armor Development
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AN OVERVIEW OF CERAMIC ARMOR APPLICATIONS
William A. Gooch Jr.
U.S. Army Research Laboratory
Weapons and Materials Research Directorate
Aberdeen Proving Ground, MD 21005-5066
ABSTRACT
The increasing capability of modern anti-armor threats and the need to field
lower weight combat vehicles, capable of engaging an opponent with little
preparation, have intensified the need for highly effective passive armor systems.
Ceramic armor technology offers significant advantages for meeting future
protection requirements, particularly for the U.S. Army’s Future Combat System.
The investigation and application of ceramics against small arms threats has a
long history, dating back to the early 1960s and the ballistic performance of
ceramic armors for personnel protection is very high; the principles governing
these defeat mechanisms and the design parameters against such threats are now
generally understood. However, achieving similar ceramic performance versus
larger caliber, kinetic energy penetrator threats have long presented a difficult
challenge. This paper presents an overview and discussion of the ballistic
requirements, ceramic design factors and a chronology of significant U.S.
developments and applications of ceramics for armor.
INTRODUCTION
The application of ceramics for armor continues to be primarily used in
lightweight armor systems for protection against small arms and machine gun
threats. The design of these systems is typically based upon the mechanical
properties of the ceramic to fracture the penetrator and the ability of a rear
compliant layer to catch the projectile debris and the damaged ceramic material.
For defeat of these low-velocity, short projectiles, the fracture mechanism occurs
very early in the process with the majority of the interaction time dedicated to
energy conversion of the kinetic energy of the debris into deformation and
delamination of compliant backing. For medium caliber and heavy armor
applications, where the dominant threat is modern, high velocity, heavy metal
eroding projectiles, the defeat mechanisms are much more complicated and of
longer time duration. For the past three decades, a wide variety of research
Ceramic Armor Materials by Design 3
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programs, both domestic and foreign, have focused on developing improved
ceramic armor systems for the defeat of these threats. This paper presents an
overview and discussion of the ballistic requirements, ceramic design factors and
a chronology of significant developments and applications of ceramics for armor,
with emphasis on research conducted on ceramic armors at the U.S. Army
Research Laboratory (ARL).
TERMINAL BALLISTIC EFFECTS
A review of the difference in terminal ballistic effects observed during the
interaction of different classes and caliber’s of kinetic energy (KE) projectiles is
important to understand the required defeat mechanisms and armor designs. The
delineation between the threat projectiles is primary related to the caliber, velocity
and energy available, but is not exact and some projectiles cross over into the two
categories discussed below. While the penetrator/target interactions for these two
categories involve similar processes, defeat of the higher performance, long rod
threats require different emphasis in the armor design parameters to be successful
and the progress has been much slower.
Small Arms/Heavy Machine Gun Defeat
Historically, ceramic composite armor systems were designed to defeat
armorpiercing (AP), kinetic energy projectiles, mainly in the small arms and
heavy machine gun category. These AP projectiles are purely inertial rounds,
most commonly made of hard steel (HRc 60-64), of moderate density (7.85
g/cm3) with a few select rounds employing even harder tungsten carbide (WC)
cores at higher densities (13.5-15.0 g/cm3). The hard core is generally encased in
a thin jacket of a more ductile metal for interior ballistic or aerodynamic
considerations, but penetration performance of the bullet is controlled by the core
properties. Such projectiles typically have a length to diameter (L/D) ratio in the
range of 3:1 to 5:1 with moderate muzzle velocities of less than 1 km/s. The
generally accepted high-end caliber is 14.5-mm, typified by the Soviet KPV
family of heavy machine guns. Some saboted, light armor-piercing (SLAP)
rounds have velocities up to 1.3 km/s but with reduced core weight. Overall, these
projectiles tend to produce a total KE on the order of 103 - 10
4 J.
Early Research [1–4] discovered that the perforation of ceramic armor
systems occurred in three general stages: 1. shattering; 2. erosion; and 3. catching.
During the shattering phase, the penetrator fractures and breaks on the surface of
the ceramic plate; the high compressive strength of the ceramic overmatches the
loading produced by the penetrator impact, and the penetrator material flows and
shatters. This initial stage is followed by a period of damage accumulation in the
ceramic material initiated by tensile wave reflections, and bending of the ceramic
4 Ceramic Armor Materials by Design
tile and backing plate. During the second stage of ceramic armor penetration, the
ceramic material is cracking, but the ceramic material can still contribute to defeat
of the penetrator core through erosion mechanisms. In the final catching phase,
the ceramic has lost considerable strength, but ceramic and backing combine to
reduce the velocity through momentum transfer mechanisms.
The defeat mechanism for hard-core AP projectiles is primarily stages 1 and 3
with projectile fracture upon impact against an armor plate having sufficient
hardness and/or high obliquity. The shattering and subsequent dispersion of the
fragments result in a dissipation of the kinetic energy of the core over a larger
area than if intact, thereby achieving defeat of the round with a reduced amount of
armor plate. Monolithic ceramic plates were best suited to produce the shattering
phenomena due to their high hardness and low densities. However, ceramic armor
requires a backup component to support the ceramic and delay failure during the
initial impact/shattering interaction; the backup component then serves to absorb
the residual projectile fragments and comminuted ceramic particles (Phase 3). The
state of the art in protection against small arms threats is typified in lightweight,
two-component ceramic faced composite armors designed for use in breast plates
for personnel body armor, armored helicopter seats and appliques to metal or
composite based vehicle structures.
Heavy Metal Long Rod KE Projectile Defeat
The mechanism for defeat of long rod penetrators (LRP) is more complex than
for the conventional AP projectiles described above. These penetrators are
commonly made of high strength, high density materials, such as tungsten
sintered alloy or depleted uranium, having densities near 18 g/cm3 with moderate
hardness, good toughness and ductility; hence, the projectiles are not susceptible
to shattering as hard core, relatively brittle, AP projectiles. This category includes
APDS and armor-piercing, discarding-sabot, fin stabilized (APDSFS) projectiles,
in calibers from 20-mm up to >140-mm. These LRPs are designed with a high
L/D ratio (currently fielded examples exceed 30:1) and the high density core
material coupled with relatively high muzzle velocity (1.3 - >1.6 km/s), yields KE
in excess of 106 J, creating a high energy density per unit area of target impacted
than with a corresponding hard core AP round. These factors, when combined
with the greater projectile length and reduced propensity for fracture, makes the
LRP a much more effective penetrator. Even if the frontal portion of the LRP can
be effectively damaged, a substantial portion of the rod remains to continue the
armor penetration process. Thus, the conditions that allow a simple ceramic
Ceramic Armor Materials by Design 5
composite to function effectively for small arms defeat do not apply when the
armor is impacted by a LRP. The primary defeat mechanism is erosion (Phase 2)
and the effectiveness is relatively low for simple ceramic armor systems.
SIGNIFICANT DEVELOPMENTS IN CERAMIC ARMOR TECHNOLOGY
Figure 1. Abex-Norton
Ceramic/ Metal Composite
Attempts to increase performance of ceramic tiles continued during the 1980’s
to present, as penetrator threats evolved. Researchers realized that increased
efficiency of ceramics might be possible by lengthening the duration of the
shattering stage of the penetration process, and/or by increasing the efficiency of
the erosion process of the comminuted ceramic material. These researchers found
that modest lateral confinement allowed constraint
of the broken ceramic pieces, thereby enhancing the
erosion phase of the ceramic penetration process.
This confinement could be obtained by casting, as
seen in the then very efficient armor developed in
1984-86 by ABEX-NORTON where silicon carbide
tiles were inserted into very accurately cast
aluminum matrices [5](Figure 1). Additional
examples include test geometry’s proposed by
Woolsey and others [6,7] to provide a stiff and
substantial confinement of the ceramic tiles in
depth of penetration (DOP) configurations.
The most significant observations during this period, however, were in 1987
by Hauver et al [8] who examined test geometry’s that delayed the generation of
damage in the ceramic tile, thereby increasing the duration of the shattering phase
of the penetrator defeat process. As penetrator threats increased in length and L/D
ratio, Hauver realized that the shattering stage duration was critical to the overall
efficiency of the ceramic defeat process; he demonstrated ceramic tile
confinement geometry’s that substantially
increased the shattering/erosion phase of the
penetrator defeat process to completely erode
the penetrator (Dwell). These experiments
employed compressive confinement of the
ceramic tile (heat shrink of the metal
surround), in combination with techniques to
delay tensile wave and bending damage to the
ceramic. The ceramic performance was
enhanced through control of system geometry
to minimize damage and increase the
shattering stage of penetrator defeat.Figure 2. Hauver’s Observation
of Ceramic Dwell in Laboratory
6 Ceramic Armor Materials by Design
However, the overall mass and space efficiencies of these laboratory packages
were low, due to the considerable confinement materials employed in the
geometry.
In 1994, research lead to the demonstration of a set of medium caliber and
full-scale armor targets that incorporated existing ceramic defeat knowledge into
an armor technology known as tandem ceramic armor (TCA) [9]. TCA
determines the optimum performance of a specific cross-sectional ceramic armor
design and then repeats the designs in multiple, shock-isolated sections; the
performance is thus additive (Figure 3). Laboratory targets, utilizing conventional
laminated ceramic-metal technology, demonstrated system designs that produced
the state of the art for KE performance. A limiting factor, however, was the space
requirements that grew as the penetrator performance increased.
TANDEM ARMOR SYSTEM
1. CERAMIC TILE
2. CONFINEMENT FRAME
3. POLYMERIC ADHESIVE
4. SUPPORTING PLATE
METAL/COMPOSITE)
5. THIN GRP SECTION
(OPTIONAL)
6. HONEYCOMB/ISOLATION
MATERIAL
7. VEHICLE HULL
Figure 3. Tandem Ceramic Armor Concept
The latest efforts to generate increased efficiency in ceramic armors are to
enhance both the erosion and “dwell” mechanisms of ceramic armor for
penetrator defeat. The development of hot-isostaticpress (HIP) processing of
ceramics with metal surrounds (Figure 4) has demonstrated dwell on the ceramic
front surface of laboratory scale
threats at efficient armor system
areal densities [10]. This HIP
processing forms a macro
composite through the
generation of residual
compressive stresses (mismatch
of thermal expansion
coefficients of the ceramic tiles
and metal confining plates) in Figure 4. Hot-Isostatic Pressed Metal
Encapsulated Ceramic
Ceramic Armor Materials by Design 7
the ceramic tile during cool down of the HIP assembly from the pressing
temperature. The macrocomposite is then able to withstand the large ballistic
bending loads during round impact, so that the ceramic tile resists fracture and
retains a high compressive strength. The macro-composite formed by HIP
processing also keeps the broken ceramic pieces confined during the second
erosive phase of the ceramic armor defeat process, should it occur, thus
maintaining a high erosive efficiency.
CERAMIC ARMOR DESIGN REQUIREMENTS
As with all armor systems, many design factors and production decisions
influence the effectiveness of ceramic armors. These processes have to be
understood and controlled to maintain performance.
Ceramic Type
The technical ceramics available for armor are numerous, but generally are
divided into the lower cost sintered and the higher cost hot-pressed ceramics. The
higher cost ceramics are justified when the lowest areal weight system is the main
requirement with the prime ceramics being boron carbide for body armor and
airborne platforms and silicon carbide for ground vehicles. The high density
tungsten carbide ceramic has specific applications where space is a limiting factor
[11]. The lower cost sintered 99.5% aluminum oxide or silicon carbide or the
reaction-bonded ceramics can be used were weight is not the driving requirement.
However, for many armor applications, ultra-high hard steels, titanium or
laminates of these materials with aluminum/composite backings are very
competitive in performance with significant engineering advantages. Table 1 lists
some of the primary producers of ceramics used today in armor applications.
Table 1. Current Ceramic Armor Producers
Ceramic Type Producer/Type
Sintered Coors CAP3 99.5% Alumina
Morgan Matroc (UK) Alumina
ETEC Alumina
Ceradyne Sintered SiC
Pure Carbon SiC
Reaction-Bonded M-Cubed (Simula) SiC
MC2 (Australia) SiC, B4C
Metal Matrix Composite Lanxide Dimox AS109
Lanxide Dimox-HT
Hot Pressed Cercom B4C, TiB2, WC
Ceradyne B4C, TiB2, SiC
Saint-Gobain B4C
8 Ceramic Armor Materials by Design
Bonding/Impedance Effects: The use of hard face ceramic materials, bonded
unto metal and composite backings, is typified in the classic work by Wilkens et
al [12,13,14] on understanding the fundamental penetration mechanics that occur
during the interaction of a hardcore steel projectile with a hard face aluminum
oxide ceramic on aluminum. The primary applications involve bonding the
ceramic to the metal or composite backing with lowdensity, low-impedance, and
low shear strength adhesives. The unfavorable impedance effects and induced
tensile failure across these boundaries and at the lateral boundaries of the ceramic
are well documented by Hauver [15,16,17] for eroding long rod penetrators.
A less understood, but equally important effect from the use of a low-shear
strength adhesive has been documented by a number of researchers. In 1993,
Furlong et al [18] presented an exact solution for the transmission of spherical
waves across planar surfaces; the coefficients of reflection and refraction were
shown to depend not only on the acoustic impedance’s of the media, but also on
the boundary conditions at the interface, the wave face curvature, and the source
frequency. Three types of boundaries can exist: 1) free, as with a free standing
ceramic plate; 2) no shear-coupling, as with two unbonded or lightly bonded
plates or 3) shear-coupling, where good adhesion or coupling exists, allowing the
transmission of transverse motion and stress. The latter shear-coupled or no slip
condition provides the best interface for a ceramic/metal armor design. Similar
investigations were conducted by Alme [19]. Leighton et al [20] discussed the
increased ballistic performance of laminated ceramic-titanium composites that
resulted from increased interlayer bond strength (strong, shear-coupled
metallurgical bonds). These effects are inherent in functionally gradient materials
(FGM) composites as observed by Gooch et al [21] where metal layers transition
into the ceramic layers without interfaces.
Adhesive Thickness and Uniformity: In simple bonded ceramic-metal
laminates, an important factor to eliminate variability in ballistic performance is
to maintain a uniform adhesive bonding layer at the minimum thickness. Burkins
[22] modified the standard DOP test configuration by examining the ballistic
results of a set of Taguchi experiments where the rear ceramic/metal bond
thickness and lateral side confinement bond thickness were varied. The least
variance occurred with a minimum bond thickness for the side and rear. For DOP
tests, the maximum bond thickness allowed for the rear and sides is 0.127-mm
(0.005 in). The uniformity of the bond thickness is maintained by placing spacers
in the adhesive.
Ceramic Armor Materials by Design 9
Confinement and Stiffness: The design of efficient ceramic systems begins by
considering the mechanisms by which a ceramic tile fails during loading and
designing the armor system to reduce the stresses contributing to early failure of
the ceramic tile. Consideration of the ballistic event with emphasis on penetrator
interface defeat on the ceramic front surface (Figure 5) lead Horwath [23] to
determine two primary areas of concern: (1) the compressional loading of the
ceramic directly under the penetrator rod, and (2) the maximum flexure of the
ceramic plate and tensile stress/strain at the ceramic plate rear surface. These two
factors are heavily influenced by the side and rear confinement thickness and
materials.
Compressive Loading Under
Projectile
Deflection and Back Surface
Stress of Ceramic tile Under
Ballistic Load
Figure 5. Primary Areas of Ceramic Tile
Multi-hit Requirements, Edge and Joint Impacts: As with all armor systems,
the requirement to provide full protection against multiple impacts is still valid for
ceramic armor designs. This requirement significantly impacts the design and tile
size of ceramic designs. Table 2 lists the U.S. Army minimum impact spacing
requirements for metals, metal laminates and ceramic laminates.
Table 2. Multi-hit Impact Requirements for Vehicular Armor
WEAPON CALIBER METALS AND METAL
LAMINATES* (mm)
CERAMIC
LAMINATES* (mm)
7.62-mm 27 54
12.7-mm 45 90
14.5-mm 51 102
20-mm 70 140
23- to 25-mm 75 145
30- to 50-mm 152 76 (105)** 152
* Minimum spacing between impacts measured from center to center of impacts
in vertical plane
** Full bore AP bullets only
10 Ceramic Armor Materials by Design
The requirement to impact edges or joints and maintain the same ballistic
performance as center tile impacts is a major design requirement for ceramic
armor systems. Generally, ceramic design is driven by the edge or joint protection
with the center providing greater protection. In some ceramic designs, the edges
of the ceramic tiles are raised to increase the thickness to equalize the protection
across the ceramic tile face. These factors result in increased areal weight for the
design.
APPLICATIONS OF CERAMIC ARMOR FOR COMBAT VEHICLES
The application of ceramics as the main protection technology has made
major advances in the last decade and represents the accepted technology, in use
today, for small arms and heavy machine gun protection, primarily as a ceramic
laminate applique over metal structural base armors and a few, newer composite
based systems. A few systems were designed against 30-mm APDS, but few
armored systems have been designed against larger threats. The following
paragraphs describe some of the military armor applications in use or
development today. These are representative, but not inclusive, of the myriad
examples of ceramic armors under development worldwide. The information was
provided by the fabricators and producers of the ceramics and products.
Armorworks
Armorworks Incorporated of Phoenix, AZ fabricates a wide range of ceramic
composite products. Shown in Figure 6 is armored kit for an AH-60H helicopter
floor. This armor system is an aluminum oxide based armor system that provides
7.62-mm APM2 protection at muzzle velocity. The armor kit consists of five
panels, two of which are removable in flight (cargo hook access) and are nested in
aft panel. The armor kit mounts on top of the floor panels using exiting fastener
points on the floor with coverage of about 5.1-m2 (55-ft
2). The armor panels
passed MIL-STD-810E environmental
testing including high and low
temperature, solar radiation, sand and
dust, salt fog, high pressure wash,
humidity, fungus, vibration-resonance
and vibration-endurance tests. The tile
and backing are bonded; the gross
panel shape is then fabricated by
cutting and grinding and diamond
saws and cores drills the holes to the
final panel configuration.
Courtesy Amorworks
Figure 6. Armorworks AH-60H Floor
Armor
Ceramic Armor Materials by Design 11
Ceradyne Incorporated
Ceradyne Incorporated of Costa Mesa, CA develops and produces a wide
range of advanced ceramics for many applications including ballistic grades such
as hot-pressed boron carbide,
silicon carbide and titanium
diboride, pressureless sintered
silicon carbide and reaction-bonded
and sintered silicon nitride.
Ceradyne has a long history of
armor development beginning in the
1960’s with the first applications of
boron carbide for combat helicopter
protection. Today, Ceradyne
designs, develops and manufactures
ceramic armor such as the ceramic
breast plates and Cobra helicopter
bucket seat of Figure 7.
Photos Courtesy of Ceradyne
Figure 7. Ceradyne Body Armor and
Helicopter Seat
Cercom Incorporated
Cercom Incorporated of Vista, CA has been a prime producer of a wide range
of commercial and ballistic grades of ceramics since 1985. Using their pressure-
assisted densification (PAD) process, Cercom has hot-pressed large quantities of
aluminum nitride, boron carbide,
silicon carbide, silicon nitride,
titanium diboride and tungsten
carbide ballistic ceramics for the
U.S. Army. Figure 8 shows Cercom
ceramic tiles on the European Tiger
helicopter seat and two different
types of Cercom hotpressed boron
carbide body armor inserts, a
single-piece, compound curvature
plate that is used in the U.S. Army
Small Arms Protective Insert
(SAPI) vest and two examples of
multiple tiles fabricated into single
protective inserts.
Photos Courtesy of Cercom
Figure 8. Cercom Ceramic Tiles for (L)
Tiger Helicopter and (R) body armor
inserts
12 Ceramic Armor Materials by Design
German Ingenieurbüro Deisenroth
The German Company Ingenieurbüro Deisenroth (IBD) of Lohmar, Germany
has established itself as a world leader in the variety and quantity of vehicles
incorporating the MEXASTM ceramic/metal/composite design; MEXASTM stands
for Modular, EXpandable Armor Systems and is composed of layered appliques
that can be added to a basic vehicle structure to give the desired protection. While
not conceptionally different from other appliques, the early use and continued
application of this design is noteworthy and at least 39 different vehicles in ten
countries utilize MEXASTM, including Austria, Switzerland, Germany, Canada,
U.S., France, Italy, Finland, Sweden and Norway. Figure 9 provides a collage of
the different vehicle types utilizing the MEXASTM system, from engineer vehicles,
tactical trucks, and numerous wheeled and tracked combat vehicles.
Figure 9. Examples of tactical and combat vehicles
Photos courtesy of NDHQ Canada and Ingenieurbüro Deisenroth
that mount IBD MEXASTM armor
The design of the MEXASTM system is shown in Figure 10 where a second
layer of protection is being placed over the first. The vehicle structure provides
the base protection and this system could be configured against three different
missions or selective uparmoring of the vehicle. The panels are mounted by
threaded attachment studs that accept special recessed fasteners.
Ceramic Armor Materials by Design 13
Photos courtesy of Ingenieurbüro Deisenroth
Figure 10. Multiple layer concept of IBD MEXASTM armor
Detroit Diesel General Motors of Canada
The Canadian National Defense Forces have been very active in providing
increased protection for a wide range of tactical and support Canadian equipment.
This requirement is driven by the deployment of their forces in a number of
peace-keeping operations and the threat of increased small arms threats. Shown in
Figure 11 is the Canadian LAV III Armored Personnel Carrier (APC) that has
protection against small arms AP threats. The ceramic MEXASTM composite
armor is fabricated by the Canadian company DEW Engineering and
Development Limited of Miramichi, New Brunswick, Canada under license to
IBD. The characteristic mounting hardware of IBD armor is readily visible in the
LAV III glacis area. DEW has supplied over 750 kits to the Canadian Defense
Forces.
Photos Courtesy of Program Manager Brigade Combat Team
Figure 11. The MEXASTM
armor panels mounted on the Canadian LAV III APC
14 Ceramic Armor Materials by Design
General Motors Defense Systems
The U.S. Army has initiated a major development program to transform the
existing family of heavy vehicles to a lighter, more agile and deployable force.
The Future Combat System (FCS) is planned for fielding by 2015. As part of the
transformation, a contract to purchase an interim family of light vehicles under
the Interim Brigade Combat Team has been awarded to GM GDLS Defense
Group L.L.C. of Sterling Heights, MI [24]. Among the many variants is the
Infantry Combat Vehicle (ICV) shown in Figure 12. Based on the LAV III chassis
and hull, the ICV mounts a version of the MEXASTM system of IBD. The
similarities in the design and mounting are visible.
Photos Courtesy of Program Manager Brigade Combat Team
Figure 12. The ICV of the Interim Brigade Combat Team utilizes
the IBD MEXASTM
applique
The GM GDLS contract indicates the ICV is to have overhead and all around
protection for the squad and crew from 152-mm Artillery high explosive airburst
at an undisclosed distance from and above the vehicle. The ICV shall also provide
integral 360 and overhead squad and crew protection from 7.62-mm AP threats
and 360 squad and crew protection from 14.5-mm AP ammunition, both fired
from undisclosed impact conditions. The ICV shall also provide the capability to
mount add on armor packages to protect against hand held shaped charge
warheads up to and including the RPG-7.
Textron Marine and Land Systems
Textron Marine and Land Systems of New Orleans, LA is the prime fabricator
for two interesting applications of ceramic composites, the U.S. Army Armored
Security Vehicle (ASV) and the Marine Corp Landing Craft, Air Cushion
(LCAC) vehicle (Figure 13). On the initial vehicle procurement, the ceramic
composite armor kit on the ASV was produced by Simula Safety Systems of
Phoenix, AZ, based on a MEXAS. license from IBD. Textron is currently
working on a new composite armor design. The ASV offers front, rear and side
Ceramic Armor Materials by Design 15
protection from 0.50-caliber armor-piercing ammunition. The LCAC is a high-
speed, over-the-beach fully amphibious, landing craft capable of carrying a 60-75
ton payload. Critical areas of the vehicle including the turbine housings are also
protected with a Simula-developed, aluminum oxide composite.
Photos Courtesy Textron
Figure 13. The Textron ASV and LCAC vehicles both mount composite armors
General Dynamics Land Systems
General Dynamics Land Systems Division (GDLS), Sterling Heights, MI has
licensed and acquired an advanced, lightweight armor technology, named
SURMAX™ Armor. This armor technology is used on the sides and rear of the
hull and the sides of the turret of the U.S. Marine Corps' Advanced Amphibious
Assault Vehicle (AAAV) to protect the vehicle from 14.5-mm AP threats and
artillery fragments (Figure 14).
Figure 14. Marine Corp AAAV and SURMAX™ being mounted
on AAAV spaceframe
Photos Courtesy General Dynamics Land Systems
SURMAX™ consists of a ceramic composite front panel attached to an armor
backing. The backing can be a composite material (such as Kevlar or S-glass) or
the structure of a vehicle (such as aluminum, steel, or titanium). The combination
of the front panel and the backing are used to stop the penetrator and the
application of SURMAX™ on the spaceframe structure of the AAAV is shown in
Figure 14. SURMAX™ is also used on the U.S. Army's wheeled Armored
16 Ceramic Armor Materials by Design
Ground Mobility System (AGMS) and the flexible panels can be fit to curved
surfaces such as wheel wells (Figure 15).
Photos Courtesy General Dynamics Land Systems
Figure 15. SURMAX™ on AGMS with curved panels in wheel wells
Extensive testing of AAAV armor at tight multi-hit distances was required for
engineering development and
Government validation tests. The
AAAV SURMAX™ side armor panel
on the left of Figure 16 is a typical
validation target, after four partial
penetrations, three with 14.5-mm AP
and one with a simulated artillery
fragment. Shots 1 and 2 are located
101-mm (4”) apart. Shown on the right
is a multi-hit test panel with ten 0.50-
caliber AP impacts for the AGMS, all
partial penetrations with the tenth shot
at a distance of 76-mm (3”) from one
previous shot.
Photos Courtesy General Land Systems
Figure 16. Ballistic Multi-hit
Qualification Tests for the AAAV and
the AGMS
Simula Incorporated
Simula Incorporated, Phoenix, AZ has been a designer and fabricator of
ceramic composite armor components since 1970 for a wide range of products
from helicopter seats, aircraft armors, body armors and ceramic composite
appliques for a range of vehicles. Figure 17 shows some of Simula’s ceramic
components. On the left is an AH64 helicopter seat that is fabricated from hot-
pressed Cercom boron carbide on Kevlar backing; the middle picture shows the
Interceptor body armor vest with hot pressed ceramic tile inserts; and the photo on
the right shows one-piece sintered silicon carbide plates made by M-Cubed of
Monroe, CT which also can be used in the Interceptor vest.
Ceramic Armor Materials by Design 17
Figure 17. Simula products: (L) AH64 Helicopter seat. (M)
Ceramic Inserts for the Interceptor Body Armor vest. (R)
Sintered silicon carbide one-piece inserts for body armor
Photos Courtesy Simula
United Defense Limited Partnership
The Ground Systems Division of United Defense Limited Partnership
(UDLP), headquartered in York, PA is one of the largest ground vehicle
fabricator’s in the U.S. The use of ceramic composite materials and structures has
been in development for many years and UDLP has progressed through three
generations of systems. The first generation development was the Composite
Infantry Combat Vehicle technology
demonstrator (Figure 18) which
replaced most of the aluminum
structure of the M2 Bradley Fighting
Vehicle with a S-2 glass reinforced
composite. This allowed the
demonstration of bonding of titanium
diboride tiles to the hull sides for
14.5-mm protection. Tile spacing
and cutouts are visible.
Photo Courtesy UDLP
Figure 18. 1st Generation CIFV with
TiB2 tiles
The 2nd
generation UDLP
ceramic composite armors can be
seen in the well-designed ballistic
protection of the M8 Armored Gun
System (Figure 19). The M8 is fitted
with bolt-on appliques and boxes
that can provide different levels of
protection from KE penetrators to
hand-held shaped charge warheads.
The ceramic composite shows
multiple impacts of 7.62-mm AP
projectiles on the test panel.
Photo Courtesy UDLP
Figure 19. M8 passive/shaped charge
armors being tested
18 Ceramic Armor Materials by Design
The 3rd
generation UDLP armor is seen in the Composite Armored Vehicle
(CAV) technology demonstrator (Figure 20). The CAV incorporates full spectrum
protection into the vehicle design, including 7.62-mm hull protection as well as
enhanced 30-mm protection to the crew station. The multi-hit performance of the
hex tiles used in the hull design as well as the excellent multi-hit protection of the
silicon carbide/titanium crew station armor against 30-mm threats is shown.
Photos Courtesy UDLP
Figure 20. CAV Technology demonstrator and hull and crew station
protection tests
CONCLUSIONS
This paper has presented an abbreviated overview and discussion of the
ballistic requirements, ceramic design factors and chronology of significant U.S.
developments over the last 30 years. The applications of ceramics for armor are
growing rapidly as the need for lighter and more agile combat vehicles increases.
Ceramic armor technology offers the best potential for meeting future protection
requirements, particularly for the U.S. Army’s Future Combat System.
REFERENCES1.
C. Donaldson, “The Development of a Theory for the Design of Lightweight
Armor”, Aeronautical Research Associates of Princeton, Inc., Technical Report
AFFDL-TR-77-114.2.
A. L. Florence, “Interaction of Projectiles and Composite Armor”, Stanford
Research Institute, AMMRC-CR-69-15, August 1969.3.
A. M. Prior, “The Penetration of Composite Armor by Small Arms
Ammunition”, Proceedings of the International Ballistic Symposium, 1986. 4.
A. K. Wong and I. Berman, “Lightweight Ceramic Armor - A Review”, Army
Materials and Mechanics Research Center, Report No AMMRC-MS-71-1, 1971. 5.
Final Report, “Demonstration of Cast, Composite Ceramic Armor (C3A), BRL
Contract DAAA-15-86-C-0014, 1990. 6.
P. Woosley, “Ceramic Materials Screening by Residual Penetration Ballistic
Testing”, 13th
International Symposium on Ballistics, June 1992.
Ceramic Armor Materials by Design 19
7. B. Morris and C. Anderson, “The Ballistic Performance of Confined Ceramic
Tiles”, 1991 Ground Vehicle Survivability Symposium, April 15, 1991. 8. G. Hauver and J. Dehn, “Interface Defeat Mechanisms in Delayed
Penetration”, 14th DEA-G-1060 Armor/Anti-armor Workshop. 9. W. Gooch, J. Prifti, P. Woolsey, J. Mackiewicz and W. Perciballi, “Tandem
Ceramic Armor for Defeat of Kinetic Energy Penetrators, ARL-TR-1946, May
1999.10.
E. Horwath and W. Bruchey, “The Ballistic Behavior of HIP Encapsulated
Ceramic Tiles”, 8th Annual TARDEC Ground Vehicle Survivability Symposium,
Monterey, CA, March 1997. 11.
W. Gooch and M. Burkins, “Ballistic Development Of U.S. High Density
Tungsten Carbide Ceramics”, Dymat 2000, Krakow, Poland, 23-29 September
2000.12.
M. Wilkins, C. Honodel and D. Sawle, "An Approach to the Study of Light
Armor", UCRL-50284, June 1967. 13.
M. Wilkins, C. Honodel and D. Sawle, "Second Progress Report of Light
Armor", UCRL-50349, Nov 1967. 14.
M. Wilkins, C. Honodel and D. Sawle, "Third Progress Report of Light Armor
Program", UCRL-50460, July 1968. 15.
G. Hauver, P. Netherwood, R. Benck, W. Gooch, W. Perciballi and M.
Burkins, "Variations of Target Resistance During Long-rod Penetration into
Ceramics", 13th Int. Ballistics Symposium, Stockholm, Sweden, 1992. 16.
G. Hauver, P. Netherwood, R. Benck and L. Kecskes,"Ballistic Performance
of Ceramics", U.S. Army Symposium on Mechanics, Plymouth, MA, 17-19
August 1993. 17.
G. Hauver, P. Netherwood, R. Benck and L. Kecskes, "Enhanced Ballistic
Performance of Ceramics", 19th Army Science Conference, Orlando, FL, 20-24
June 1994. 18.
J. Furlong, C. Westbury and E. Phillips, “A Method for Predicting the
Reflection and Refraction of Spherical Waves across Planar Interfaces”, J. of
Applied Physics, Vol. 76, July 1994. 19.
M. Alme, Alme Associates, private communication. 20.
K. Leighton, R. Franz, and A. Gerk, “Laminated Ceramic-Titanium
Macrocomposite Armor”, 8th Annual Ground Vehicle Survivability Symposium,
Monterey, CA, 24-27 March 1997. 21.
W. Gooch, M. Burkins and R. Palicka, “Development And Ballistic Testing Of
A Functionally Gradient Ceramic/Metal Applique”, NATO Applied Vehicle
Technology Panel, Loen, Norway, 7-11 May 2001. 22.
M. Burkins and W. Gooch, “U.S. Ceramic Ballistic Test Methodology and
Data”, TTCPWTP1 Meeting, Maribyrnong, Australia, 10 May 1995.
20 Ceramic Armor Materials by Design
23. W. Bruchey and E. Horwath, “System Considerations Concerning the
Development of High Efficiency Ceramic Armors”, 17th Int. Sym. on Ballistics,
Midrand, South Africa, March 1998. 24.
DoD Contract DAAE07-00-D-M051, 16 November 2000, Brigade Combat
Team web site.
Ceramic Armor Materials by Design 21
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ARMOR CERAMICS UNDER HIGH-VELOCITY IMPACT OF A
MEDIUM-CALIBER LONG-ROD PENETRATOR
Hans-Jürgen Ernst, Volker Wiesner and Thomas Wolf
French-German Research Institute of Saint-Louis (ISL)
P.O. Box 34, F-68301 SAINT-LOUIS CEDEX (France)
ABSTRACT
In the first part of this paper, continuous measurements of the cratering
process in unconfined targets (Al2O3) of different lateral dimensions as well as in
completely confined constant-volume targets (Al2O3, B4C, SiC, and TiB2) are
given; they are achieved with a penetration gauge developed at ISL. The
confinement and material-related influences on the penetration resistance are
shown.
Secondly, the protective power of these ceramics is quantified and compared
to that of other inert materials. By means of an exponential fitting function, which
is based on the assumption that the very beginning of the penetration process is
not influenced by the geometrical armor configuration, appropriate ballistic
material parameters, called ductile limit of the space equivalence factor, are given.
Based on this parameter, a ballistic screening of materials is presented, which
enables a configuration-independent comparison of the protective power of the
investigated ceramics with that of other inert materials. In conclusion, it shows
that brittle materials are still interesting for light-weight armor design.
INTRODUCTION
The ballistic performance of a single ceramic material depends on both the
thickness of the ceramic block and its constructive environment, often called
confinement1. It is known that the ballistic performance of thick targets decreases
for most of the ceramic materials with increasing thickness 2, 3, 4
. The more a
ceramic material is confined, the more it tends to behave like a ductile one,
achieving a higher ballistic performance 5.
Time-dependent measurements of penetration into steel-confined ceramic
targets are not easily realized, as material thickness and relatively small density
differences exceed the capabilities of common measuring techniques (i.e. X-rays).
Ceramic Armor Materials by Design 23
To the extent authorized under the laws of the United States of America, all copyright interests in this publication are the propertyof The American Ceramic Society. Any duplication, reproduction, or republication of this publication or any part thereof, withoutthe express written consent of The American Ceramic Society or fee paid to the Copyright Clearance Center, is prohibited.
That is why, a special terminal ballistic gauge for continuous penetration
measurements has been developed at ISL6, 7
.
This paper has two aims: firstly, a closer insight into the time-dependent
penetration behavior of some ceramic materials is given, based on new results
with the penetration gauge. Secondly, a simplified choice of brittle materials for
light-weight armor design is proposed by introducing a ballistic material
parameter that allows a
configuration independent
evaluation of the protective power
of inert armor materials.
EXPERIMENTAL BASIS
The upper left-hand part of
figure 1 shows a photograph of the
kinetic energy projectile BMU G
154 developed at ISL. Below a
sectional drawing of the rod is
presented. Some penetrator
material data are given in the upper
right-hand table. In the lower left-
hand diagram the reference
penetration Pref is plotted versus
the impact velocity vz
kept constant in this
paper8: at 1800 m/s an
RHA penetration of
about 165 mm is
achieved. The Pref(vz)-
function and its fitting
parameters a and b can
be found in the lower
right-hand corner.
Figure 1: KE projectile BMU G 154
As seen in figure 2,
ceramic blocks of
different thicknesses
(tile thickness: 20 mm,
lateral dimensions:
100x100 mm2;
materials: Al2O3, B4C,
SiC and TiB2) have
been investigated inFigure 2: target set-up and ceramic data
24 Ceramic Armor Materials by Design
three configurations: "unconfined", “laterally” and "totally confined". Some
ceramic material data are found in the table. In every experiment the ceramic
block thickness tz, and the residual penetration Pres are measured; the reference
penetration Pref is evaluated with the Pref(vz)-function.
Continuous penetration measurements can be made advantageously by means
of a special gauge developed at ISL7. Figure 3 shows in its upper left-hand part a
schematic drawing and below a sketch explaining the working mechanism of this
gauge. It consists of a metallic tube with an electrical resistance wire inside placed
in a hole in the target material (Ø approx. 1 mm) along the expected penetration
axis. The projectile dynamically closes the contact between tube and wire. The
ongoing penetration process causes an electrical resistance decrease that is
measured as a time dependent tension variation.
The upper right-hand part of figure 3 shows a photograph of the gauge; the
evaluation of the gauge signals is explained in the lower right-hand picture. Apart
from the transient beginning of the penetration process and its final phase (from
ceramics into RHA) the experimental penetration curve is approached by a
polynomial P(t) fit function, which yields the cratering velocity u(t) after
differentiation.
Figure 3: ISL penetration gauge
Ceramic Armor Materials by Design 25
PROTECTIVE POWER DEFINITIONS
Based on experimental results (see figure 2), the total penetration Ptot , the
average density tot of the test set-up, as well as an RHA layer thickness tref ,
which is equivalent to the ceramic block, can be derived 9:
Ptot = tz + Pres , tot = (tz · z = Pres · ref ) / Ptot and tref = Pref – Pres . (1)
Normalized formulations for ceramic thickness tz,n , residual penetration Pres,n and
target density z,n
tz,n = tz / Pref , Pres,n = Pres / Pref and z,n = z / ref , (2)
as well as for total penetration Ptot,n , reference layer thickness tref,n and total
density tot,n
Ptot,n = tz,n + Pres,n , tref,n = 1 – Pres,n and tot,n = tot / ref , (3)
will mostly eliminate the experimentally caused scattering.
Though the protective power of a target set-up against a defined threat can
directly be quantified by the total penetration, normalized ballistic factors are
advantageously used for this purpose. Here, equivalence factors F describe the
volume gain (subscript s) and the mass gain (subscript m) of the ceramic block
under consideration, as compared to the equivalent reference material layer:
Fs = (1 – Pres,n ) / tz,n and Fm = Fs / z,n . (4)
In order to complete the formulations, efficiency factors E are added, which
describe the volume gain (subscript s) and the mass gain (subscript m) of the total
penetration in the test set-up Ptot , as compared to the reference penetration Pref :
Es = 1 / (tz,n + Pres,n ) and Em = Es / tot,n . (5)
Based on the assumption that the very beginning of the penetration process is
not influenced by the geometrical armor configuration, an appropriate ballistic
material parameter, called ductile limit of the space equivalence factor, has been
introduced 10
. The exponential fitting function for the space equivalence factor Fs
and the appropriate formulation for the normalized residual penetration Pres,n are:
Fs = Fs (0) · exp( ·tz,n ) and Pres,n = 1 – Fs (0) ·tz,n · exp( · tz,n) (6)
26 Ceramic Armor Materials by Design
Fs(0) is the
configuration-
independent ballistic
material parameter. It
quantifies the space
equivalence that is
reached, if the ceramic
material behaves during
penetration like a ductile
one. The approximation
coefficient depends on
both the material and the target configuration.
Figure 4: thickness-dependent residual penetrationand space equivalence of inert materials
Figure 4 shows qualitative diagrams of the proposed approximation function
Fs (left-hand side) and of the resulting Pres,n (right-hand side) dealing with a
hypothetical brittle, ductile and composite material, in order to describe different
types of penetration behavior schematically.
RESULTS
Figure 5: confinement influence on the ballistic
performance of brittle materials
Confinement Influence
When a projectile hits an
unconfined ceramic target, the
ceramic material in front of the
projectile-target interaction zone
is increasingly fractured by the
shock wave and its reflections11
. Depending on the lateral tile
dimensions, the predisturbed
ceramic material can expand
radially, thus causing a density
decrease in the material.
Sufficient lateral dimensions
and/or a well-designed
confinement are able to reduce
or even to stop this expansion
completely5.
The upper diagram of figure
5 firstly shows the decrease in
residual penetration behind 120-
mm thick unconfined blocks of
glass and Al2O3 as a function of
the lateral tile dimension a.
Ceramic Armor Materials by Design 27
Secondly, a decrease in residual penetration can be seen for laterally confined
blocks (a=100 mm). By following the horizontal dotted lines, significantly higher
lateral tile dimensions for the unconfined configuration are found.
The lower left-hand diagram of figure 5 shows a comparison of gauge
measurements for totally confined and unconfined Al2O3 blocks of equal
thickness. Appropriate second-order polynomial fits of the experimental curves
are given in the table. In the lower right-hand diagram the corresponding cratering
velocity comparison can be seen. The confined configuration shows the lowest
slope of the penetration curve and a significantly smaller cratering velocity, thus
indicating a higher penetration resistance as compared to the unconfined
configuration.
Influence of the Ceramic Material
Further time-dependent measurements were made in order to compare the
penetration resistance of 120-mm-thick totally confined blocks of some often used
armor ceramics (Al2O3,
B4C, SiC and TiB2). In
the left-hand diagram
of figure 6 penetration-
versus time plots are
given for these
ceramics and,
additionally, one for
mild steel. The
equations of the
second-order
polynomial fits
corresponding to the
measured curves are
presented in the table
below. The right-hand
diagram shows the
corresponding cratering
velocities as a function of time. Though the transient region of the penetration
beginning is not well defined, it is obvious that the penetration curves of all
ceramics have a significantly lower slope than that of mild steel, signifying a
higher penetration resistance. SiC and TiB2 have the highest curvature.
Figure 6: penetration gauge measurements in differentceramic materials
Differentiation of the P(t)-polynomials yields straight lines that indicate the
average cratering velocities. The u(t)-curves of all ceramic materials start at lower
values, and with the exception of Al2O3, they have higher slopes than that of mild
28 Ceramic Armor Materials by Design
steel. Two - eventually accumulating - types of ceramic penetration behavior may
explain these effects:
1. the lower the cratering velocity at the beginning of the penetration process,
the more the ceramic reacts like a rigid target (examples: TiB2 and SiC);
2. the higher the slope of the u(t)-curve the more the projectile is decelerated
during the ceramic penetration process (example: B4C).
Protective Power of Some Ceramics
Detailed results of the DOP experiments with TiB2 and B4C, Al2O3 and SiC are
presented in the four upper diagrams of figure 7. Each of them shows
experimentally determined Fs-data and Fs(tz,n)-curves calculated by using equ. (6)
for the three
investigated
configurations. In
the lower part of
the figure the
protective power
hierarchy of inert
materials is
presented in the
form of a
diagram, in
which the ductile
limit of the space
equivalence is
plotted against
the normalized
density.
Figure 7: protective power of some brittle materials
It can be seen
that TiB2 has the
highest Fs(0)-
value of the
investigated
ceramics; a better
confinement may
still increase its
penetration
resistance. The
space equivalence
factors of Al2O3
Ceramic Armor Materials by Design 29
are similar to those of B4C. The latter is more interesting for armor use because of
its significantly lower density. In the case of
SiC, the Fs(0)-values range between those of TiB2 and B4C.
In the lower diagram it can be observed that the Fs(0)-value of TiB2 is higher
than those of high-hardness steels, the latter being followed by that of SiC. No
significant differences exist between the ductile limits of the aluminas, B4C,
Si3N4, and titanium; that of glass, aluminium and GFRP materials share a lower
position. It can be observed that the (generally used) Fm(0)-values of high-
hardness steels and GFRP materials are comparable, but their Fs(0)-values
(signifying the ballistic result) differ significantly. Obviously, the good Fm(0)-
values of B4C and of the GFRPs are mainly due to their low densities. Though
GFRPs additionally profit from their increasing thickness-dependent space
equivalence 12
, it can be concluded that ceramics are still useful as light-weight
armor materials.
CONCLUSIONS
Continuous gauge measurements offer a closer insight into the penetration
process of ceramics impacted by a LRP at 1800 m/s. A heavy confinement cuts
the initial cratering velocity in Al203 down like a rigid target and it increases the
deceleration of the projectile. These two effects may explain the protective power
of different armor ceramics too: TiB2 and SiC have relatively low initial cratering
velocities; B4C decelerates the projectile due to a continuously decreasing u(t)-
curve. By optimally using the material intrinsic and/or the structural confinement,
the quasi-ductile penetration behavior of some ceramics might be approached.
A diagram, in which the ductile limit is plotted against the normalized density,
is perhaps more useful to quantify the protective power of ceramics (and other
inert materials) than the solely used mass factor. This graph shows on the one
hand that well-confined ceramics have comparable (SiC) or even higher ductile
limits (TiB2) than high-hardness armor steels and on the other hand it also shows
that the good protective power of B4C is mostly due to its relatively small density.
REFERENCES 1
2
3
Westerling L., Lundberg T., "The Influence of Confinement on the
Protective Capability of Ceramic Armour at Two Different Velocities", 15th ISB,
Jerusalem, Israel, 1995
Andersen Jr. C.E., Walker J.D., Lankford J., "Investigations of the Ballistic
Response of Brittle Materials", SWRI-Technical Report, 1995
Yaziv D., Partom Y., "The Ballistic Efficiency of Thick Alumina Targets
against Long-Rod Penetrators", 14th ISB, Quebec, Canada, 1993, Vol. 2, pp. 331-
340
30 Ceramic Armor Materials by Design
4
5
6
7
8
9
10
11
12
Hauver G.E., Netherwood P.H., Benck A.F., Gooch W.A., Perciballi W.J.,
Burkins M.S., "Variation of Target Resistance During Long-Rod Penetration into
Ceramics", 13th ISB, Stockholm, Sweden, 1992
Ernst H.-J., Hoog K., Wiesner V., "Ballistic Impact Behavior of Some
Ceramics in Different Environments", EURODYMAT 94, Oxford, UK, 1994
Wiesner V., “Erfassung der Projektilbewegung im Ziel mit
Widerstandssonden“, MEBAL 85, ISL R 116/85, 1985
Ernst H.-J., Hoog K., Wiesner V., Wolf T., “DOP and Continuous Cratering
Measurements in Differently Confined Ceramics”, EAFV Symp., Shrivenham,
UK, 1996
Rapacki E.J., Hauver G.E., Netherwood P.H., Benck R.F., “Ceramics for
Armors – a Material System Perspective”, 7th TARDEC Ground Vehicle Symp.,
USA, 1996
Ernst, H.-J., Merkel Th., “Zur Vereinheitlichung der Anwendung
ballistischer Faktoren“, ISL RT 519/2000, 2000
Hoog K., Ernst H.-J., Wolf T., “A New Parameter Characterizing the
Ballistic Performance of Ceramics”, EURODYMAT 97, Toledo, Spain, 1997
Bless S.J, Subramanian R., Partom Y., Lynch N., “Effects of Radial
Confinement on the Penetration Resistance of Thick Ceramic Tiles”, 15th ISB,
Jerusalem, Israel, 1995
Ernst H.-J., Wolf T., Unckenbold W., “Protective Power of Thick
Composite Layers Against Medium-Caliber Long-Rod Penetrators”, 19th ISB,
Interlaken, Switzerland, 2001
Ceramic Armor Materials by Design 31
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PRACTICAL ISSUES IN CERAMIC ARMOUR DESIGN
Dr Bryn James
Defence Science and Technology Laboratories
Chobham Lane
Chertsey, Surrey, KT16 0EE
United Kingdom
ABSTRACT
The performance of ceramic armour is heavily dependent upon the
configuration of the system. Generally, compromises must be made in factors
such as single-shot performance in order to obtain the best overall system
performance and to accommodate practical requirements such as multi-hit
capability.
This paper will discuss some of the factors involved in designing practical
armour systems and will use experimental results to illustrate some design
techniques and improvements that may be made. Factors to be addressed include
the physics of stress propagation across an interface, acoustic impedance
matching, optimisation of ceramic tile edge profile, general rules on ceramic
armour design and choice of ceramic material.
INTRODUCTION
Ceramic materials are capable of displaying significantly better protective
performance than an equivalent weight of metal armour. The ability to perform so
well depends partly upon the very strong dependency, inherent in all ceramics, of
the yield stress on ambient pressure (1). It is therefore apparent that, for an armour
system to display the performance intrinsically available from the ceramic
elements, the armour configuration must be such that the ambient pressure is
maintained at the highest levels possible. In addition, the ambient compressive
stress should be highly homogeneous, in order to avoid stress gradients giving rise
to shear and tensile stresses which can lead to early catastrophic failure of the
system.
It is possible to devise an experimental system for mechanically constraining
ceramic materials so that the required conditions are met. Several examples of
these confinement systems are reported in the literature [2, 3, 4, 5, 6]. Such
Ceramic Armor Materials by Design 33
To the extent authorized under the laws of the United States of America, all copyright interests in this publication are the propertyof The American Ceramic Society. Any duplication, reproduction, or republication of this publication or any part thereof, withoutthe express written consent of The American Ceramic Society or fee paid to the Copyright Clearance Center, is prohibited.
confinement systems may be capable of maintaining a uniform compressive state
within the ceramic so that the highest levels of strength may be obtained resulting
in complete erosion of the penetrator on the surface of the ceramic tile. This
performance would be of great benefit in an armour system. However, the
confinement system required to maintain such a stress state generally has
significant mass and relies upon accurate impact of the projectile on a
predetermined site. These two factors indicate that the commonly used,
experimental confinement systems are inappropriate for practical ceramic armour.
Instead, to maximise the performance of a practical ceramic armour system, it is
necessary to attempt to mimic the effects of massive confinement by suitable
manipulation of the stresses generated during impact, using the lightest possible
configuration.
The requirements of practical ceramic armour systems generally go beyond
the maximisation of single impact performance. Often, an additional requirement
is that the system should be capable of defeating several impacts within a given
area, i.e. it must have a multi-hit capability.
IMPEDANCE MATCHING
The compressive wave launched from the site of a projectile impact
propagates as a crude approximation to a spherical wave (in the far field). As this
compressive wave impinges upon any interface defined by a change in acoustic
impedance, a number of transmitted and reflected waves are produced. In general,
any interface between two dissimilar materials will give rise to a set of both shear
and longitudinal reflected and transmitted waves, even if there is a very close
match in acoustic impedance.
Figure 1. Stress transformation at an interface
34 Ceramic Armor Materials by Design
Figure 2. Shear failure of adhesive bond between alumina and aluminium
Metal backed ceramic armour is often constructed using a polymer adhesive.
In such cases, any advantage of close impedance matching of the metal and
ceramic layers is lost due to the inclusion of the low impedance glue layer. The
change of impedance at this layer gives rise to a strong tensile reflection into the
ceramic which acts to shatter the ceramic layer and to provide energy for ejection
of the shattered material. In addition to this, a strong shear wave is set up at the
interface which serves to ‘unzip’ the adhesive interface. An example of such shear
‘unzipping’ can be seen in Figure 2, where concentric bands of shear failure of the
toughened epoxy adhesive layer can be seen.
Very few adhesive materials exist with an acoustic impedance close to that of
metals or ceramics. The class of adhesives with the closest acoustic match are the
high temperature use ceramic adhesives. Typically these may have an impedance
of 6 MRayls compared to 37 MRayls for alumina and 46 MRayls for RHA steel,
reducing the stress reflected back into the ceramic by 33%. However, even this
relative mismatch is considerably better than that for polymeric materials for
which the stress reflected back into the ceramic may be 98%. Unfortunately, such
ceramic adhesives are not as strong as polymer glues and they must therefore be
used as a matching layer in a mechanically bonded system. When such a matching
layer is used, the improvement in performance can be considerable. An example
of the difference in impact performance when a suitable matching layer is
incorporated can be seen in Figure 3.
Ceramic Armor Materials by Design 35
Figure 3. The effects of acoustic matching
Figure 3A shows a typical armour system coverplate impacted by a 30mm
calibre APFSDS projectile when no matching compound is used. The illustration
at Figure 3B shows the same coverplate when a few grams of acoustic matching
cement are used. Similarly Figure 3C shows the resultant ceramic debris when no
matching compound is used and Figure 3D shows the intact tiles (side view) after
impact when the matching compound is incorporated.
All our experience shows that an adhesive with an enhanced level of acoustic
impedance pays dividends in improving the multi-hit performance, the structural
integrity and, to a lesser extent, the ballistic mass effectiveness of composite
ceramic armour. It remains a technological challenge to produce a high strength
adhesive with such an acoustic impedance.
TILE EDGE GEOMETRY
Typical ceramic armour packs are built from arrays of tiles bonded in some
way to a backing. Due to lack of constraint, and stress wave reflections at the tile
edges, the protective capability for projectile impact at the tile edges is often
reduced. In order to provide a specified minimum level of protection, steps must
be taken to manage this performance reduction by such means as thickening the
36 Ceramic Armor Materials by Design
tile edges or using thicker tiles. Both of these approaches are costly in terms of
weight or price.
Figure 4. Edge cracking of ceramic tile due to tensile reflections
In principle, if the impact induced stress wave can be propagated efficiently
across the tile boundary, the edge integrity may be maintained for an extended
time during the impact, enhancing the edge performance.
Figure 4 indicates failure induced at the tile edge due to stress waves
propagating away from the impact site being reflected at the edge. Impact closer
to the edge exacerbates this behaviour.
We used a 7.62mm, precisely made, experimental steel projectile
(performance closely matched to 7.62mm APM2) to test the performance of a
number of edge profiles of alumina tiles glued to 8mm aluminium alloy (7017)
plates. The residual energy of the emerging projectile was measured using high
speed photography and a measurement of the residual mass.
A number of different edge profiles were investigated, including profiling the
tile thickness and the tile width. The most successful profile was shown to be a
45 chamfer on the tile edges. Results can be seen in Figure 5, where the triangles
labelled LAT 45 indicate the performance of the 45 edge chamfer system on a
6mm tile. For comparison, the performance of an 8.5mm thick tile with standard
edges can also be seen, showing that, close to the tile edge, the 8.5mm tile
protection is only as good as that in the centre of a 6mm tile. By using chamfered
edges, a weight saving of 9.2 kg m-2
or 30% of the ceramic mass was
Ceramic Armor Materials by Design 37
demonstrated, for a specified protection level, (7) without resorting to thickened
edge tiles.
Figure 5. Relative protective performance of different edge profiles
38 Ceramic Armor Materials by Design
Figure 6. Configurations used in edge optimisation evaluation
OPTIMUM CERAMIC/BACKING RATIO
For many years, armour designers have used a general rule that 2/3 of the
mass of a composite ceramic armour system should be invested in the ceramic
front layer, whilst the backing system should consist of 1/3 of the total mass. For
an alumina/aluminium system this corresponds to a ceramic/metal thickness ratio
of 1.5. Whilst this relationship gives a good first estimate for the optimum
configuration for low velocity impact at normal incidence, we have found that it
does not give an adequate estimation of the optimal ratio over a range of velocity
or for oblique impact.
There is little information available in the literature that bears direct
comparison, but analysis of the available data shows a significant correlation for a
wide variety of projectile types impacting an alumina/aluminium system.
Hetherington (8) refers to work by Ali (9) showing an optimum ceramic
thickness, in the ballistic limit configuration, for the defeat of 7.62mm AP rounds
impacting alumina/aluminium systems at ~850 ms-1
. The latter showed
experimentally that maximum ballistic limit velocity, VBL, was obtained for a
ceramic/metal plate thickness ratio, Tcer/Tmet, of 1.5 for normal impact (VBL =
850 ms-1
), reducing to 1.0 for 30 obliquity.
Hohler, Stilp and Weber (10) used a somewhat more complex target structure,
incorporating a thin RHA and rubber front layer. However, their results using an
8.2mm diameter tungsten sinter alloy rod, with an enlarged central section,
impacting at 1500 ms-1
, give optimum thickness ratios, Tcer/Tmet, of
approximately 2.0 at 0 obliquity, 1.25 at 45 obliquity, and 0.82 at 60 obliquity.
The results of Hohler et al. (11) in another study show an optimum
performance for a ceramic/metal plate thickness ratio of 0.71 for alumina on
aluminium at 60 obliquity and at 1450 ms-1
using a 71mm long, L/D=20, tungsten
alloy rod. This value is consistent with that obtained by Hetherington and by
Hohler, Stilp and Weber. The latter similarity is not surprising as the impact
conditions were similar. The similarity in the optimum Tcer/Tmet ratio found by
Hohler et al. with that of Hetherington is, however, quite surprising given the
difference in impact conditions. It can be seen in Table 1, that the optimum
Tcer/Tmet ratio is highly dependent upon the impact conditions. It can be seen
that this ratio changes to 1.7 for Al2O3/RHA at 1450ms-1
and 60 obliquity, to 4.9
for SiC/Al at 1450ms-1
and 60 obliquity and to a ratio of 15.0 for SiC/Al at 2200
ms-1
and 60 obliquity. This ratio is an indicator of the relative performance of the
ceramic and metal fractions. It can be seen that the relative performance of the
ceramic increases with impact velocity and decreases with obliquity.
We can now revise the general rule that the hard ceramic front layer should
contain 2/3 of the system mass whilst the supporting back layer contains 1/3 of
Ceramic Armor Materials by Design 39
the mass to include the effects of impact and velocity. A more useful
approximation to the optimal thickness ratio has been devised by fitting the,
admittedly sparse, available data to the following simple equation:
angle)Impact90(x000,60
Velocity)optimum(
TT
met
cer (3)
where Velocity is in ms-1
and Impact angle is in degrees.
The applicability of this fit can be seen by reference to the following table:
Table 1. Optimum thickness ratio for alumina/aluminium armour systems
Velocity
ms-1
Impact Angle
degrees
Experimental
Optimum Ratio
Tcer/Tmet
Optimum Ratio
From
Equation 3
Source
850 0 1.5 1.28 Ref. (8)
850 30 1.0 0.85 Ref. (8)
1450 60 0.71 0.73 Ref. (11)
1500 0 2.0 2.26 Ref. (10)
1500 45 1.25 1.13 Ref. (10)
1500 60 0.82 0.75 Ref. (10)
CHOICE OF CERAMIC MATERIAL
The ballistic mass effectiveness of ceramic materials is dependent upon the
armour configuration and the threat projectile. However, if the experiment is well
designed, a reliable general ranking of mass effectiveness may be measured
across a range of threats. We have performed such measurements using 14.5mm
heavy machine gun rounds and 30mm and 40mm APFSDS projectiles. Some of
the materials investigated are detailed in Table 2. An average ballistic mass
effectiveness has been calculated for impact from these projectiles in a number of
configurations, for these materials. Results can be seen in Figure 7. From this
figure it would appear that titanium diboride should be the ceramic of choice.
However, it is no secret that titanium diboride is expensive.
A cost analysis was performed using the best, informed estimate of the price
of ceramic materials in ‘production’ quantities (sufficient supply for 100 generic
light armoured vehicles). It can be seen that the hot pressed non-oxide materials
are significantly more expensive than sintered materials, but where either mass
(Figure 9) or thickness of armour (Figure 10) are critical, use of a non-oxide
40 Ceramic Armor Materials by Design
ceramic is indicated. It should be noted that this cost estimate depends upon
several factors, the most significant being the use of a given ceramic for non-
armour applications. Armour is generally a relatively small sales area for a
ceramic production company. The cost of production of a material can only be
reduced if large quantities are required for another, non-armour application.
Identification of a bulk-use application for a specific material could result in a
dramatic reduction in the price of a ceramic material for armour.
Table 2. Ceramic materials used in cost/mass/thickness analysis
Name Material type
RHA United Kingdom RHA HV30 = 3.39 Gpa 7840 kg m-3
Alumina 1 Sintered 95% alumina 3680 kg m-3
Alumina 2 Sintered 98% alumina 3780 kg m-3
Novel alumina DSTL developed novel sintered alumina 3690 kg m-3
RB-SiC Reaction bonded silicon carbide 3210 kg m-3
TiB2 Hot pressed titanium diboride 4520 kg m-3
B4C Hot pressed boron carbide 2520 kg m-3
SiC Hot pressed silicon carbide 3230 kg m-3
AlN Hot pressed aluminium nitride 3270 kg m-3
When we calculate the mass of material required to defeat a given threat
(Figure 9), there is surprisingly little variation across a wide range of materials. It
can be seen that we pay a lot more for a small increase in performance. If we
calculate a figure of merit (1/(cost x mass2) (as mass is more important than cost
in our application) for all of these materials (Figure 11), it can be seen that
alumina becomes the most attractive material.
Figure 7. Ballistic Mass Effectiveness
Ceramic Armor Materials by Design 41
Figure 8. Total cost of material required to defeat a given threat
Figure 9. Mass required to defeat a given threat
Figure 10. Thickness required to defeat a given threat
42 Ceramic Armor Materials by Design
In Figure 11, we see that “Novel alumina” has the highest figure of merit of
all the materials studied. This material, a sintered alumina with modified
microstructure, was developed within the DSTL Armour Physics Group as part of
a programme to study the effects on performance of changing the microstructure
of alumina. It can be seen that it is possible to make significant improvements to
the performance of alumina to improve its attractiveness as an armour material. It
is believed that yet further improvements are possible.
Figure 11. Figure-of-merit considering cost and mass
ACKNOWLEDGEMENT
The work upon which this analysis is based was funded by the UK
Government Corporate Research Programme.
I would like to thank my colleagues Antony Barker, of DSTL, and Christian
LeGallic, of DGA, France, for their contribution to the experimental work.
REFERENCES
1T. J. Holmquist, G. R. Johnson, W.H. Cook, “A computational constitutive
model for concrete subjected to large strains, high strain rates and high
pressures”, 14th International Symposium on Ballistics. Sept. 1993 2
S.J.Bless, M.Benyami, L.S.Apgar and D.Eylon, “Impenetrable ceramic targets
struck by high velocity tungsten long rods”, Structures under shock and impact
II, Ed. P.S.Bulson, Computational Mechanics Publications, June 1992.3
G.E.Hauver, P.H.Netherwood, R.F.Benck and L.J. Kecskes, “Ballistic
performance of ceramic targets”, Proc. Army Symposium on Solid Mechanics,
Plymouth, Mass. USA, Aug 1993.
Ceramic Armor Materials by Design 43
4 P. Lundberg, R Renstrom and L. Holmberg, “An experimental investigation of
interface defeat at extended interaction time”, Proc. 19th
International
Symposium on Ballistics, pp. 1463-1470, May 2001. 5 B.James, "The influence of the material properties of alumina on ballistic
performance", 15th International Symposium on Ballistics, May 1995. 6 N.S.Brar, H.D.Espinosa, G.Yuan and P.D.Zavattieri, “Experimental study of
interface defeat in confined ceramic targets”, Proc. APS Topical conference on
shock compression of condensed matter, July 1997. 7
GB Patent Application Number 0026710.4, “Ceramic Tile Armour”,
26th
October 2000, B. James 8 J. G. Hetherington, Two component composite armours, Proc. Light Weight
Armour Systems Symp., Shrivenham, UK, (1995) 9 M. S. B. Ali, “Optimisation of composite armour for normal and oblique
impact”, MSc Thesis, 21 Military Vehicle Technology Course, RMCS,
Shrivenham, UK, (1993) 10
V. Hohler, A. J. Stilp and K. Weber, “Ranking methods of ceramics and
experimental optimization of a laminated target with ceramics”. Proc. Light
Weight Armour Systems Symp., Shrivenham, UK, (1995) 11
V.Hohler, K. Weber, R. Tham, B. James, A.Barker, I. Pickup, “Comparative
analysis of oblique impact on ceramic composite systems”, Proc. Hyper
Velocity Impact Symposium, Nov. 2000
44 Ceramic Armor Materials by Design
BALLISTIC DEVELOPMENT OF TUNGSTEN CARBIDE CERAMICS
FOR ARMOR APPLICATIONS
Dr Pierre-François Peron
Etablissement Technique de Bourges
Route de Guerry
18021 Bourges Cedex
France
ABSTRACT
In the frame of a cooperative research project agreement, France and the
United States of America are developing and optimising ballistically a new class
of ceramics which offers a very high space effectiveness. These ceramics have a
higher density than armor steel (about 15) and should solve protection weaknesses
on vehicles due to space restrictions. In this paper, the elaboration process and the
mechanical characteristics of these “high density” ceramics are first detailed.
Their ballistic performances against 44 APFSDS medium caliber projectile are
then presented.
INTRODUCTION
A cooperative research project agreement has been signed in December 1996
between the Minister of Defense of the French Republic and the Secretary of
Defense of the United States of America as regards the study of “high density”
ceramic technology for armor applications. The aim was to develop and to
optimise a new class of ceramic which offers a good mass effectiveness and
mainly a high space efficiency. These ceramics should solve vehicle protection
weaknesses related to space restrictions. They have a greater density than Rolled
Homogeneous Armor (RHA) steel and are designed for applications on medium
armor and high armor vehicles.
France and the United States have independently developed high density
ceramic belonging to tungsten carbide ceramics (WC) which density is about 15.
The French WC materials are WC/metal cermets with low metal binder content
while the U.S. WC materials are high purity WC with no binder addition.
During the cooperation, dynamic properties of these two kinds of ceramics are
investigated and their ballistic performances are evaluated against kinetic energy
Ceramic Armor Materials by Design 45
To the extent authorized under the laws of the United States of America, all copyright interests in this publication are the propertyof The American Ceramic Society. Any duplication, reproduction, or republication of this publication or any part thereof, withoutthe express written consent of The American Ceramic Society or fee paid to the Copyright Clearance Center, is prohibited.
projectiles and shaped charges to optimise target parameters and to compare the
results with other ceramics. Some material exchanges are also carried out to
enlarge the range of threats.
The French research is conducted by the Etablissement Technique de Bourges
(ETBS), Bourges, France, and the Centre Technique d’Arcueil (CTA), Paris,
France. The U.S. research is conducted at the Weapons and Material Research
Directorate of the U.S. Army Research Laboratory (ARL), Aberdeen Proving
Ground, MD. This paper documents the development of the French ceramics and
provides some ballistic test results against 44 APFSDS kinetic energy threat.
HIGH DENSITY CERAMICS
The ceramic class designated as “high density” includes all the ceramics
which density is greater than that of RHA steel (7.85). A review of potentially
interesting ceramics was undertaken by CTA1 and showed that a large number of
ceramic oxides, nitrides, carbides and borides met the density criteria. Mechanical
properties of some of them are listed in table I.
However, most of them were difficult to process industrially or had prohibited
costs. Tungsten carbide ceramics exhibited high mechanical properties and had a
great deal of applications on the civilian and the military markets. This kind of
ceramic was thus chosen for further investigations.
Table I. Physical and mechanical properties of various high density ceramics.
Ceramic Density
Melting
point
(OC)
Modulus
(GPa)
Hardness
(kg/mm2)
Poisson
ratio
Longitudinal
wave
velocity
(m/s)
MoB 8.77 2600 400 2350 -- 6754
TaC 14.5 4170 285 1800 0.24 4810
TaN 14.3 3093 -- -- -- --
UO2 11 3140 193 6800 0.3 4450
WB 16 2900 -- -- -- --
WC 15.8 3050 660 2700 0.22 6900
WC-Co 14.95 1450 645 2600 0.22 6900
46 Ceramic Armor Materials by Design
SANDVIK TUNGSTEN CARBIDE CERMETS
The French material is a tungsten carbide ceramic with a low content of
cobalt. The cobalt is used as a binder to process WC cermet at lower pressure and
temperature as compared to pure WC ceramic. During the cermet elaboration, the
binder is added as a liquid phase to the WC powder. The agglomerated powder is
then densified by iso-static compression and sintered between 1350°C and
1500°C.
The French WC cermets are produced by Sandvik Hard Material society
located in Epinouze, France. They are available in 200 mm square or 250 mm
diameter cylinder tiles and in thickness up to 40 mm.
Simulations conducted with the analytical code BREFIL2 showed that the best
ballistic performances should be reached with WC-Co cermets which had less
than 10 % in weight of binder and a fine microstructure. Three cermets were thus
chosen for ballisitic evaluations : H3F, H5F and H10F. Their physical and quasi-
static properties are provided in table II. These cermets have a cobalt content
between 3 % and 10 % in weight and a very fine grain size (less than 1 m). The
addition of cobalt entails a decrease of the ceramic hardness and of its
compressive strength but increases its flexural strength.
Table II. Physical and mechanical properties of Sandvik and Cercom WC ceramic.
H3F H5F H10FWC
Cercom
WC content
(% weight) 97 92.5 89.5
96.8% WC
2.8% W2C
Co / other contents
(% weight) 3 / 0 4.5 / 3 10 / 0.5 0 / 0.4
Grain size
Average size Extra fine
Extra fine
0.5 µm Extra fine
Extra fine
0.9 µm
Theoretical density
Measured density
15.3 15.
14.95
14.5 15.7
15.6
Melting Point (°C) 1450 2800
Hardness (HV30) 1925 1810 1600 2200
Flexural Strength (MPa) 2570 2640 3960 1100
Tenacity (MPa m) 7 9 13 7
Young modulus (GPa) 670 645 580 690
Poisson ratio 0.21 0.21 0.22 0.20
Sound speed (m/s)
Longitudinal
Transversal
6970
4222
6858
4300
Ceramic Armor Materials by Design 47
CERCOM TUNGSTEN CARBIDE CERAMIC
The U.S. material is a pure tungsten carbide ceramic. Its elaboration process
has been developed by the Cercom incorporated of Vista, CA in USA and enables
to produce high purity WC ceramic with no binder. The Cercom WC ceramic
contains 96.8% WC and 2.8% W2C in weight and has a density of 15.6. Its
physical and mechanical properties are listed in the table II and compared to the
Sandwik cermets ones.
TEST CONFIGURATION
French ballistic evaluations of the Sandvik and Cercom materials were carried
out with the 44 APFSDS kinetic energy projectile produced by GIAT Industries
(Figure I).
Figure I. 44 APFSDS kinetic energy projectile.
233
The 44 APFSDS is a L/D 25
tungsten alloy rod laboratory
penetrator (Table III) which
represents the 105 APFSDS
projectile at third scale. Its has a
9.3 mm equivalent diameter and
weights 0.257 kg. Its baseline
penetration into RHA steel is 160
mm for an impact velocity of
1500 m/s.
Table III. 44 APFSDS physical and
mechanical characteristics.
44 APFSDS
Composition
(weight %) 93W-4,6Ni-2,4Fe
Density 17,6
Hardness 423 Hv30
Yield stress 980 MPa
Ultimate stress 1150 MPa
Elongation 9 %
It is nominally fired at 1500 m/s
from a 44 mm bore diameter gun
based on a 40 mm L 70 Bofors tube
and fitted with a 105 mm HM2 breech.
For the tests, the gun was put at 61 m from the target. The impact velocity was
measured by two optical barriers. The total yaw angle was calculated from the
trace of the fins in paraffin cardboard put at regular intervals close to the target.
The tests with a yaw impact superior to 1° were disregarded.
48 Ceramic Armor Materials by Design
The ballistic tests were
conducted by using the Depth of
Penetration (DOP) technique
described in figure II. This
technique compares the
performances of a RHA steel semi-
infinite target to the residual
penetration obtained in a RHA
steel block put behind a ceramic
tile. More details on the test
configuration are given in the next
section.
RHA st
f
Figure II. Target configuration
in the DOP test.
Steel lateral
confinementRHA
WC ceramic
Tcer Pres
eel
ront confinement
The ballistic performances of the ceramic are calculated from the residual
penetration of the projectile in the RHA steel back-up and are represented by an
equivalent thickness (Eeq), an equivalent mass (Meq) and a quality factor (q2) :
where PRHA is the projectile penetration in a RHA steel
semi infinite target,cer
resRHA
eqT
PPE
cer
RHAeqeq EM
eqeq
2 EMq
Pres is the residual penetration of the projectile in
the RHA steel block behind the ceramic tile,
Tcer is the ceramic thickness,
RHA and cer are respectively RHA steel and
ceramic densities.
In the case of a RHA steel plate confinement put in front of the ceramic, the
thickness of this plate is added to the residual penetration (Pres).
The reference penetration is that obtained with the 44 APFSDS at 1500 m/s in
a semi-infinite RHA steel target. It corresponds to an equivalent thickness and an
equivalent mass equal to 1. Higher values indicate that the tested material has
better ballistic performances than RHA steel.
WC SANDVIK BALLISTIC PERFORMANCES : FIRST EVALUATIONS.
A first evaluation was carried out with the three Sandvik cermets in order to
determine the influence of the cobalt content on the ballistic performances.
The tests were conducted with cylindrical cermets tiles of 240 mm diameter and
30 mm thickness. The target was composed of a ceramic module and a RHA
semi-infinite back-up (Figure III). The module was constituted of two ceramic
tiles put between two RHA steel plates of 10 mm thickness and confined laterally
by 10 mm of mild steel. All the material surfaces in contact with the ceramic were
grounded to obtain a perfect contact between the steel plates and the tiles.
Ceramic Armor Materials by Design 49
Figure III. Test configuration for the ballistic evaluation of Sandvik cermets
against 44 APFSDS.
10 10 40
RHA
RHA
Mild steel
PVC
WC ceramic
44 APFSDS
1500 m/s
The results of the ballistic evaluations are provided in the table IV and are
compared to those obtained with Al2O3 and SiC ceramics.
Table IV. Ballistic performances of several ceramics against 44 APFSDS.
Ceramic
diameter
(mm)
Ceramic
thickness
(mm)
Ceramic
surfacic
density
(kg/m2)
Impact
velocity
(m/s)
Residual
penetration
(mm)
Meq Eeq q2
RHA steel 0 0 1500 160 1.00 1.00 1.00
H3F ( 240) 2 x 30.5 933.3 1478 80.2 0.67 1.31 0.88
H5F ( 240)2 x 29.8
2 x 29.8
891
891
1503
1503
70.1
61.0
0.79
0.87
1.51
1.66
1.19
1.45
H10F ( 240) 2 x 30.3 877.2 1488 69.2 0.79 1.50 1.22
Al2O3 (94%) 50 181 1469 115.5 1.80 0.89 1.6
SiC 49.9 157.2 1469 119.2 2.04 0.81 1.66
This first evaluation performed with H3F, H5F and H10F proved that a WC
cermet with cobalt content of 5 % in weight has the most interesting ballistic
performances. The equivalent thickness is always superior to 1, which shows the
interest of these cermets in term of space. Their equivalent mass is quite lower
than 1. But the thickness of the cermet was not optimised and better performances
should be obtained with finer tiles.
Besides, because of their high density, the WC cermets ballistic performances
are specific and opposed to SiC and Al2O3 ones which are interesting in term of
equivalent mass.
50 Ceramic Armor Materials by Design
FURTHER EVALUATIONS OF SANDVIK AND CERCOM WC CERAMICS
PERFORMANCES AGAINST 44 APFSDS.
The influence of the elaboration process was also studied by comparing the
ballistic performances of H5F and WC Cercom ceramics. The test configuration
was the same as before. The results of the evaluation are listed in the table V.
Table V. Ballistic performances of H5F and WC Cercom ceramics
against 44 APFSDS.
Ceramic
diameter
(mm)
Ceramic
thickness
(mm)
Ceramic
surfacic
density
(kg/m2)
Impact
velocity
(m/s)
Residual
penetration
(mm)
Meq Eeq q2
RHA steel 0 0 1500 160 1.00 1.00 1.00
H5F 250
(cover plate)
2 x 30
2 x 30
897
897
1491
1502
72.9
76.6
0.76
0.73
1.45
1.39
1.1
1.01
Cercom 200
(cover plate)
2 x 30.1
30.2+30.1
939
941
1495
1502
38.9
57.2
1.01
0.86
2.02
1.71
2.04
1.46
These tests confirmed the ballistic performance level of H5F. The WC
Cercom ceramic performances were scattered but were higher than H5F ones.
Besides, the observation of the targets after shot (Figure IV) gave us some
explanations on the performance differences between Sandvik and Cercom
ceramics and on the WC ceramic dynamic behaviour during the penetration of the
projectile.
For firing tests on H5F cermets, the crater in the cover plate had a diameter
slightly superior to the projectile caliber. The cermet was damaged in the area of
interaction with the projectile and seemed to have been slightly affected
elsewhere. For shots on WC Cercom ceramics, the hole in the front plate was far
wider than the projectile diameter and the ceramic was damaged in a large area
around the projectile penetration zone.
According to these observations, the H5F cermet seems to exhibit limited
resistance to the projectile penetration. The binder agglomerates the WC grains
and enables to get good quasi-static mechanical properties. But at high strain
rates, it may be a weak area in comparison to WC grains and thus a privileged
way for the damage propagation in the cermet.
For Cercom ceramic, WC grains are intimately linked between each other and
exhibit thus a higher resistance to the projectile penetration. In these conditions, a
larger area of the ceramic takes part in the projectile erosion. This leads to a wide
damaged area but also an ejection of the affected ceramic through the cover plate.
Ceramic Armor Materials by Design 51
Figure IV. Interaction of 44 APFSDS with H5F Sandvik and WC Cercom ceramics.
Crater in the cover plate and damage caused to the ceramic.
Cover plate
WC ceramic damage
No Image.
Pres 72.9 mm Pres 38.9 mm Pres 57.2 mm
H5F WC Cercom WC Cercom
CONCLUSION
In the frame of a cooperative research program, France and the United States
are developing and optimising a new class of ceramic which have a high space
effectiveness. These materials are tungsten carbide ceramics and have a density
superior to that of RHA steel. They are designed for applications on medium
armor and high armor vehicles in areas where space restrictions are present. The
French materials are WC/metal cermets with low metal binder content while the
U.S. are high purity WC with no binder addition.
In France, these ceramics were evaluated in 60 mm thickness against
44 APFSDS projectile. The shots showed that both ceramics exhibit high space
effectiveness. WC Cercom and H5F ceramic performed an equivalent thickness
superior to 1.5 and an equivalent mass slightly inferior to 1. Higher ballistic
effectiveness should be obtained by reducing the ceramic thickness.
REFERENCES1 C. Cottenot, “State of art and evaluation of high density ceramics as armor
materials”, ETCA 93 R 153 (1993). 2 S. Fouquet, “BREFIL : an analytical model for the interaction between a
kinetic energy projectile and a brittle material”, ETCA 88 R 042 (1988).
52 Ceramic Armor Materials by Design
BALLISTIC DEVELOPMENT OF U.S. HIGH DENSITY TUNGSTEN CARBIDE CERAMICS
William A. Gooch and Matthew S. Burkins U.S. Army Research Laboratory Weapons and Materials Research Directorate Aberdeen Proving Ground, MD 21005-5066, USA
Richard PalickaCercom Incorporated, 1960 Watson Way Vista, CA, 92083, U.S.A.
ABSTRACT The United States and France, under a cooperative research agreement have
developed a new class of high density ceramics which inherently provide high space efficiency and reduced susceptibility to damage accumulation effects in thick sections. While many ceramics were considered, this research has focused on tungsten carbide based ceramics. The U.S. Army Research Laboratory, in cooperation with Cercom Inc. has developed a hot-pressed tungsten carbide ceramic for ballistic applications. This paper will present a survey of high density ceramics, document the mechanical and elastic properties of the U.S. WC ceramic and baseline the ballistic performance.
INTRODUCTIONIn December 1995, a cooperative Project Agreement under the Memorandum
of Understanding between the Secretary of Defense of the United States of America and the Minister of Defense of the French Republic concerning Technology Research and Development Projects was signed to jointly develop and optimize a new class of ballistic ceramic materials that offer very high space effectiveness for applications where inherent space restrictions are present. These ceramics are defined as any ceramic with a density greater than that of rolled homogeneous armor (RHA) steel (7.85 g/cm3). The U.S. research is being conducted at the Weapons and Materials Research Directorate of the U.S. Army Research Laboratory (ARL), Aberdeen Proving Ground, MD and the French research is being conducted by the Établissement Technique de Bourges, Bourges, France and the Centre Technique d.Arcueil, Paris, France. This paper documents
Ceramic Armor Materials by Design 53
To the extent authorized under the laws of the United States of America, all copyright interests in this publication are the propertyof The American Ceramic Society. Any duplication, reproduction, or republication of this publication or any part thereof, withoutthe express written consent of The American Ceramic Society or fee paid to the Copyright Clearance Center, is prohibited.
the development of the U.S. ceramic as well as providing limited ballistic testing of the ceramic.
HIGH DENSITY CERAMICS While a number of ceramic oxides, nitrides and carbides meet the criteria of a
high density ceramic, most of the ceramics are difficult to process or the costs are prohibitive. A review of possible ceramics of interest was undertaken by Cercom [1] and Table I lists important properties for nominally pure ceramics whose densities are greater than 7.85 g/cm3. Of the twelve ceramics listed, only two have been tested ballistically. Limited testing of hafnium carbide (HfC) was conducted by Hauver [2] in 1/5 scale tests to study dwell. Rupert et al [3,4] with Nuclear Metals Incorporated of Concord, MA examined uranium oxide (UO2)ballistically in both sintered and hot isostatic pressed conditions. The work was successful in producing near-theoretical density UO2 in 10-mm disks of thickness of 11.3-mm, but further work was discontinued because of the associated radiation limitations imposed by these uranium-based materials. Limited testing of high metal content WC cermets were also conducted at ARL with 14.5-mm armor-piercing projectiles, but the performance was generally equal to equivalent areal weights of RHA.
The data in Table I, while compiled from laboratory test data, established the direction for further development. The tungsten carbide (WC) family was selected as the prime ceramic of interest because of the high density, excellent mechanical properties and the potential unique applications in both the military and civilian markets. The WC family had both the highest density and modulus and exhibited a longitudinal sound speed about half that of lower density ceramics, resulting in an impedance about 2.5 times that for RHA.
Table I. Elastic Properties/Melting Temperature of Selected High Density Ceramics
CERAMIC DENSITY (g/cm3)
MELTING POINT
(°C)
MODULUS (GPa)
HARDNESS(kg/mm2)
LONGITUDINAL WAVE VELOCITY
(m/s) MoB 8.77 2600 400 2350 6754Mo2C 9.18 2522 533 1499 7620NbN 8.31 2300 483 1525 7626TaB 14.19 3090 400 3130 5309TaC 14.40 3985 285 1720 4449TaN 14.36 3087 576 -- 6333HfB2 11.19 3380 500 2900 6685HfC 12.67 3890 360 3830 5330HfN 13.39 3000 500 1600 6112UO2 10.97 2850 -- -- -- WC 15.7 2800 696 2200 6600W2C 17.20 2785 420 2150 4940
54 Ceramic Armor Materials by Design
U.S. TUNGSTEN CARBIDE CERAMICSTungsten and carbon form two ceramics of interest, tungsten monocarbide
(WC) and ditungsten carbide (W2C) as seen in the two narrow phase stabilityregions at 50 and 30 atom % carbon, respectively, in the binary phase diagram ofFig.1[6]. Both ceramics have a melting temperature of about 2800°C and WC hasa very high elastic modulus. Currentlyproduced WC materials are, in fact,cermets, alloys of ceramics and metalbinders that are sintered to form a hard dense material. These cermets containeight to ten percent cobalt by weight,added as a liquid-phase sintering aid toallow the material to be fully densified atlower temperatures and pressures ascompared to binderless WC. The cobaltaddition reduces yield strength andhardness, but increases toughness. WCcermets of these high metal content exhibitreduced ballistic performance, as the resistance to penetration is governed bythe percentage of metal content andlocation of the sintering aid in themicrostructure. The specific materialproperties and structure that make WCcermets valuable industrial materialsinherently degrade their performance asballistic materials.
Figure 1. W-C Phase Diagram
The U.S. WC ceramic processing technology was developed by CercomIncorporated of Vista, CA and the physical, mechanical and elastic properties ofthese ceramics provide a ballistic response similar to high quality ballistic ceramics of lower density. Cercom developed the processing technology to densify large, high purity ceramics without metal sintering aids, and adapted the process to the densification of WC [5]. Cercom first densified tiles of 100-mm x 100-mm size in thicknesses up to 25.4-mm. The process was scaled-up to 152-mm x 152-mm tile sizes in thickness up to 50-mm and 203-mm diameter tileswere produced in thicknesses up to 30-mm. Ditungsten carbide (W2C) was also densified in 152-mm x 152-mm size tiles of 26-mm thickness, but will not bediscussed in this paper. The French ceramic is a low-metal binder WC that wasdiscussed in the preceding paper.
Ceramic Armor Materials by Design 55
The hot-pressed Cercom WC had a density of 15.6 g/cm3 and the tiles were analyzed to be composed of WC and 2.8% W2C, the latter a byproduct of the densification process. The nominal purity was 99.6% WC/W2C. The tiles were densified without metal sintering aids allowing near-theoretical density tiles to be hot-pressed in large tile sizes. The grain size was between 0.3-1.4µm with an average grain size of 0.9µm. The crystal structure of WC and W2C is hexagonal and matches other higher performing hexagonal ballistic ceramics such as -SiC,
-TiB2, AlN, and -Al2O3. The measured quasi-static mechanical and elastic properties of the Cercom WC are provided in Table II and are compared to hot-pressed silicon carbide densified by the same process.
Table II. Measured Mechanical and Elastic Properties of Cercom WC and SiC
WC SiC-B
DENSITY (g/cm3)THEORETICAL
AS-PRESSED 15.715.6
3.223.20
Average Grain Size (µm) 0.9 4.0
HARDNESS (VICKERS-1-kg) (kg/mm2) 2200 ± 20 2700
FLEXURAL STRENGTH (MPa) WEIBULL MODULUS
1100 ± 130 10.2
65518
KIC TOUHGNESS (MPam1/2)SINGLE ETCH NOTCHED BEAM
INDENTATION VICKERS 7.56 ± 0.51 6.86 ± 0.19
5.2--
TENSILE STRENGTH (MPa) 589 ± 57 592
ELASTIC MODULUS (GPa) 690.1 455
SHEAR MODULUS (Gpa) 287 195
POISSON RATIO 0.20 0.14
SONIC VELOCITY (km/s) LONGITUDINAL
TRANSVERSE 6.8584300
12.257.65
When compared to traditional low-density ceramics, the compactness of the ceramic is a direct function of the inherent densities of materials. Relative to steel, WC is one-half the thickness for the same areal density; one-fifth that of silicon carbide ceramic. The remainder of this paper will document the ballistic performance of WC for a representative tungsten long-rod penetrator.
56 Ceramic Armor Materials by Design
TEST PROJECTILEThe 162-g tungsten
projectile, shown in Figure 2, has a Length to Diameter(L/D) of 13 and has beenused for many years as a testsimulant for a mediumcaliber long-rod projectile.Table III lists thecomposition and typicalmechanical property data onthis penetrator. The baseline RHA penetration of this rodhas been documented byGooch et al [7] and is governed by the followingtwo parameter exponentialequation where P is in mm and V is in km/s:
Figure 2. L/D 13 Tungsten Rod
Table III. Mechanical Properties for L/D13 Rod
Designation L/D 13 W A Rods
Composition (wt %) 93W-3Ni-2Fe
Density (g/cm3) 17.7
Hardness 40-45 RC
Yield Strength 1200 MPa
Ultimate Tensile Strength 1280 MPa
Elongation 8 %
2)/447.1(8.308P Ve
This equation, based on work of Lanz and Odermat [8] and improved by Goochet al [7], is accurate between 500-1800 m/s. The L/D13 rod, when nominally firedat 1550 m/s, has a baseline penetration into RHA of 129.2-mm.
EXPERIMENTAL SETUP The test penetrator was fired from a laboratory gun consisting of a 40-mm
L70 Bofors breech assembly with a 38-mm smoothbore barrel. A custom-builtpolypropylene sabot system was used to launch the projectiles. The gun waspositioned 1.5 meters in front of the targets and an orthogonal flash radiographicsystem [9] was used to measure projectile velocity, pitch, and yaw. Propellantweight was adjusted to achieve the desired striking velocity and ballistic resultswith 2° total yaw were disregarded.
The ballistic test data presented in this paper were conducted using the depthof penetration (DOP) technique developed by Woolsey et al [10] and shown inFig. 3. This technique compares the performance of a RHA baseline (PO) to theresidual penetration (PR) of ballistic tests of different thicknesses of ceramic. All tests are shot at 0° obliquity and the target had no cover plate, but was confinedlaterally by a steel frame in which the ceramic tile was epoxied. The DOP into the rear RHA plate (PR) was measured for each impact. Burkins and Gooch [11],
Ceramic Armor Materials by Design 57
when examining the sources of variancein DOP data, determined that bond thickness is a major source of variabilityin ballistic data. This observation hasled to a modification in the assemblyprocedures for DOP testing at ARL. Thebond thickness is maintained at 0.5-mm for both the side confinement and the rear interface. The DOP technique iscost-effective, but only provides arelative performance indication; time-dependent effects, reported by Hauver etal[12] predominate in this methodology where no cover plate is used, interfaceeffects are not minimized, and ceramic confinement is limited. Generally, the result of these factors are a rapid reduction in the relative performance as ceramicthickness increases, accompanied with an increase in the scatter of the ballisticdata. While these factors are present in DOP testing of the lower-densitytechnical ceramics, particularly for thicker tiles, high density ceramics appear to have less scatter. The optimum ceramic thickness for maximum ballisticerformance in simple laminate target designs, such as the DOP configuration, has been observed to be between 25 to 40-mm for the low densityceramics and thicker tiles begin to loose performance as time dependentfactors predominate. This loss primarily relates to geometric considerations ofthe ceramic tile to the confinement, the sound speed of the ceramics and time forreflection from the ceramic back and side interfaces. The material has failed andmass and space effectiveness rapidly decrease as the penetrator encounters,essentially, granulated ceramic. As the RHA penetration performance of a rodincreases, this ceramic design problem becomes greater for the armor designer,as the residual penetration becomes larger as the ceramic gets thicker. As withmany armor designs, the target is driven by the space factor, not the massfactor. While Hauver et al [13] and Prifti et al [14] were successful in overcomingsome time dependent effects in ceramics, the total mass and/or space effectivenessvalues are still low if the total parasitic mass or space is included in theeffectiveness calculations. The problem resides with the low density of theceramics for these high performance applications in simple armor designs.
Figure 3. Depth of PenetrationTest Configuration
BALLISTIC CHARACTERIZATIONBallistic performance of armors or elements of armors are characterized by
dimensionless factors which compare the areal density (mass/area) and thicknessof the material to baseline RHA. Many variations and terminologies exist, butFrank [15] developed and described a concise set of mass and space effectiveness
58 Ceramic Armor Materials by Design
factors whose conventions are in use at ARL (Figure 4). Since the DOP technique determines the equivalent RHA performance of the ceramic relative to the semi-infinite penetration of the rod, the ballistic characterization of the ceramic can be defined by the mass effectiveness (em), the space effectiveness (es) and the armorquality factor (q2) as described by the equations below; the small e indicating thatthe performance indices are elemental rather than system effectiveness. The termPRHA represents how much baseline RHA penetration was removed by ceramicthickness, TCER, at the same impact velocity and is obtained by subtracting theresidual RHA penetration depth (PR) from the baseline RHA penetration of the rod (PO), i.e., (PO - PR). The ceramic mass effectiveness can then be related to eS
by the ceramic density ( CER) and the RHA density ( RHA). RHA has an em and es
of 1.0 and higher indices indicate better ballistic performance. The quality factorhas significance for armor designers as this factor relates both the mass and spacefactors; values over 1.0 indicate armors or materials which are thinner and/or lighter than the baseline RHA performance and indicate superior armors or materials.
sm
RHA
CERAMIC
m
CERAMIC
RHAs
CERAMICCERAMIC
RHARHAm exeqxe
T
Pe
xT
xPe 2
Figure 4. Mass and Space Efficiency Parameters
EXPERIMENTAL RESULTS/DISCUSSIONTable IV documents 12 DOP tests for three thicknesses of WC tiles that were
152-mm X 152-mm in lateral size. The effectiveness factors for each test havebeen determined from the impact velocity. These 12 tests represent a consistentset of data and demonstrate a number of interesting observations. First, theballistic performance for the 10-mm and 20-mm WC tiles was very similar withthe mass effectiveness near 1.55 and the space effectiveness near 3.0. The 20-mmtiles were slightly better performers. The 30-mm WC tiles demonstratedsignificantly better performance than would be predicted from the thinner tilesand an optimum thickness may not yet have been reached. However, the residualpenetration is approaching zero and a higher performing rod will have to beutilized to determine the optimum tile thickness for WC. For the 30-mm tiles, the mass effectiveness is over two and the space effectiveness is over four. The highspace factors in these tests mean that the rod is being stopped in a ceramicthickness that is about one-fourth the penetrator length.
Ceramic Armor Materials by Design 59
Table IV. L/D 13 DOP Ballistic Test Data for WC
SIZE (mm)
CERAMIC THICKNESS
(mm)
CERAMIC AREAL
DENSITY (kg/m2)
IMPACT VELOCITY
(m/s)
AVERAGE PR (mm)
em es q2
RHA 0 0 1550 129.2 1.00 1.00 1.00
152
10.310.310.210.2
160.1160.1158.1159.3
1531154315521554
95.097.099.099.0
1.531.541.501.52
3.053.022.993.01
4.684.654.494.57
152
20.120.220.120.2
314.2314.6314.2315.0
1541154515541561
62.062.069.071.0
1.651.651.521.49
3.283.293.022.96
5.415.444.614.40
152
30.230.230.030.1
471.5470.7467.6470.3
1543154315431549
7.05.87.7
11.4
2.022.042.031.97
4.014.054.033.91
8.108.268.177.68
CONCLUSIONS This paper documents the U.S. development of a new class of ballistic
ceramics known as high density ceramics under a joint U.S/France cooperative research program. These ceramics have been defined to be any ceramic whose density is greater than that of rolled homogeneous armor steel (7.85 g/cm3) and are very interesting from an armor standpoint, as very compact targets are possible. The WC family of high density ceramics was selected as the prime ceramic of interest because of the high density, excellent mechanical properties, and the potential applications in both the military and civilian markets. The tungsten carbide ceramic densification technology was developed by Cercom Incorporated of Vista, CA, who succeeded in producing high purity WC ceramics in large tile sizes. Ballistic testing with a L/D13 tungsten rod demonstrated very high space effectiveness factors.
With increasing battlefield threats, current and future combat vehicles will require armor technologies which obtain maximum protection with compact structures and armors. The development of this ceramic provides armored system developers with a very space efficient material for use against higher-performing medium caliber and full-scale rods in applications such as add-on appliques, roof appliques, hatch appliques or hull and turret side armors.
60 Ceramic Armor Materials by Design
REFERENCES 1.
2.
3.
4.
5.
6.
7.
8.
9.
10.
11.
12.
13.
14.
15.
S. Schneider, Engineered Materials Handbook, Vol.4, Ceramics and Glasses, American Society for Materials International, 1991
G. Hauver, Private Communication N. Rupert, R. Schoon, .Evaluation of High Density Ceramics for Ballistic
Applications., Conference on Dynamic Loading in Manufacturing and Service, Melbourne, Australia, 1993
N. Rupert, M. Burkins, W. Gooch, M. Walz, N. Levoy, E. Washchilla, Development of High Density Ceramic Composites for Ballistic Application., Inter. Conference on Advanced Composite Materials, Wollengong, Australia, 1993
Cercom Patents 5,302,561, 5,358,685, 5,354,536, Monolithic, Fully Dense SiC Material and End Uses.
E. Rudy, AFML-TR-65-2, Part IV, Compendium of Phase Diagram Data, pg. 192, Air Force Materials Laboratory, Wright-Patterson Air Force Base, June 1969
W. Gooch, M. Burkins, K. Frank, .Ballistic Performance of Titanium against Laboratory Penetrators., 1ST Australasian Congress on Applied Mechanics, Melbourne, Australia, 1996
W. Lanz and W. Odermat, .Penetration Limits of Conventional Large Caliber Antitank Guns/Kinetic Energy Projectiles., Proc. 13th Inter. Symposium on Ballistics, Stockholm, Sweden, 1992
C. Grabarek and L. Herr, .X-Ray Multi-Flash System for Measurement of Projectile Performance at the Target., U.S. Army Ballistic Research Laboratory Technical Note 1634, September 1966
P. Woolsey, S. Mariano, and D. Kokidko, .Alternate Test Methodology for Ballistic Performance Ranking of Armor Ceramics., 5th Annual U.S. Army Tank-Automorive Command Survivability Conference, Monterey, CA, 1989
M. Burkins and W. Gooch, .Ceramic Testing Methodology., U.S. Army Research Laboratory Workshop, June, 1995
G. Hauver, W. Gooch, P. Netherwood, R. Benck, W. Perciballi and M. Burkins, Variations of Target Resistance During Long-rod Penetration into Ceramics., Proc. 13th Inter. Symposium on Ballistics, Stockholm, Sweden, 1992
G. Hauver, P. Netherwood, R. Benck, .Ballistic Performance of Ceramic Targets., 13th Army Symposium on Solid Mechanics, Plymouth, MA, 1993
J. Prifti, P. Woolsey, W. Gooch and W. Perciballi, .Advanced Ceramic/Metallic Armor Systems for Defeat of Long Rod Penetrators., Second Ballistic Symposium on Classified Topics, Johns Hopkins University, 1993
K. Frank, .Armor-Penetrator Performance Measures., Armament Research and Development Command /Ballistic Research Laboratory Report MR-03097, 1981
Ceramic Armor Materials by Design 61
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To the extent authorized under the laws of the United States of America, all copyright interests in this publication are the propertyof The American Ceramic Society. Any duplication, reproduction, or republication of this publication or any part thereof, withoutthe express written consent of The American Ceramic Society or fee paid to the Copyright Clearance Center, is prohibited.
INITIAL TESTS ON CERAMICS IN COMPOSITE ARMOR
W. Lanz
RUAG Land Systems (formerly Swiss Ordnance Enterprise)
Allmendstrasse 86
CH-3602 Thun, Switzerland
ABSTRACT
The intended development of the new Swiss Main Battle Tank at the
beginning of the Seventies instigated a major move in terminal ballistics
research. The existing homogeneous armor steel was at its limits; therefore,
new armor materials had to be selected to ensure tank crew protection.
Moreover, the influence of different geometries on terminal ballistics was
investigated. In this research, ceramics were included as candidate materials to
protect against both shaped charges and kinetic energy projectiles. Mainly,
model tests served to select a suitable material or suitable material
combinations. The first candidate ceramic armor material was alumina,
followed by porous silicon nitride (Si3N4) and silicon carbide (SiC). The SiC
composite armor was also tested in full scale; although it did not perform as
well as in the model tests, it still demonstrated a very high stopping power.
INTRODUCTION
The Near East wars of 1967 and 1973 clearly
showed how vulnerable even heavy tanks had
become. Especially, the massed deployment of
guided long range antitank missiles had a
devastating effect (Figure 1).
The frontal armor of heavy tanks of that period
was homogeneous steel of about 250-mm
thickness, measured in the horizontal plane. On
the other hand, the opposing shaped charges had a
penetration performance of 500-mm or more
(Figure 2). Even the KE (kinetic energy)
projectiles fired from tanks penetrated around 300-
mm RHA (rolled homogeneous armor).
Figure 1. Effect of an 80-mm Shaped
Charge Impact on a 40-tonne Tank
Ceramic Armor Materials by Design 63
Any further thickness
increase of the conventional
armor would have led to
prohibitive vehicle weights and
would have been virtually
ineffective in view of the shaped
charge penetration power. In
order to protect the tank crews
from this threat a totally different
approach had to be chosen.Figure. 2. Before 1980: Antitank Projecti
Armor Protection of Heavy Tanks
EARLY ATTEMPTS
Of course, the problem described above was well known to all related
development institutions which carried out extensive research on projectile/target
interaction. The main objective was to design cost effective armor protection with
significantly higher stopping power than steel. The primary purpose was the
protection against the high performance shaped charge, but also against KE
rounds as well. These cause less penetration, but have a far higher energy content.
Basic analytic and experimental terminal ballistics investigations had been
described in 1, 2, 3 and other unclassified literature. Its principal statement was
the hydrodynamic approximation formula for a shaped charge jet penetration
behavior:
t
p
L
P
P = Penetration Depth p = Projectile Material Density
L = Projectile Length t = Target Material Density
The approximate formula contains no material strength values, since the
dynamic pressures at very high velocity impacts are an order of magnitude higher
than the material strengths. This hydrodynamic approach also forms the basis for
the well known "Odermatt Formula" for KE projectiles 4 this formula includes
the influence of material properties since these cannot be neglected after all.
Figure. 2. Before 1980: Antitank Projectiles vs.
Armor Protection of Heavy Tanks
64 Ceramic Armor Materials by Design
FIRST TERMINAL BALLISTICS INVESTIGATIONS AT RUAG LAND
SYSTEMS
At the beginning of the Seventies, the Swiss Defense Procurement Group
initiated the development of an indigenous main battle tank. This also entailed a
terminal ballistics research program at RUAG Land Systems which concentrated
on the evaluation of new armor protection materials and on the influence of
geometry on terminal ballistics.
When this program was initiated, sufficiently accurate terminal ballistics
calculations were not believed to be feasible. Thus, an experimental approach was
chosen. Our main model ammunitions are shown in Fig. 3.
1972
Cal. 40 mm Shaped Charges, ca. 90 g of HE
Pentastite P = 180 mm RHA
Octastite P = 225 mm RHA
1975APDS
Ø20 90 mm Tungsten, mP = 300 g
v 1300 m/s
P 90 mm RHA
1978
APDSFS
v 1500 m/s
2 Types (Tungsten):
Ø9 130 mm, mP 150 g, P 130 mm RHA
Ø8 170 mm, mP 150 g, P 165 mm RHA
at 100 mm
stand off
Figure 3. RUAG Model Ammunitions
NEW ARMOR MATERIALS EVALUATION
A closer look at the hydrodynamic penetration formula readily reveals a
feasible way to enhance armor stopping power without increasing weight, i.e.,
using lower density material. However, a drawback is the correspondingly higher
thickness required. This is shown in Figure 4 as the penetration depth PCM of a
given projectile type versus the density ratio of candidate material and steel
Ceramic Armor Materials by Design 65
CM/ St for different composite materials. The penetration PSt of the same
projectile in RHA steel of 825 MPa tensile strength serves as a reference value.
As defined in the Space Equivalence Factor,
St
CM
CM
StS
P
PF
describes the protection performance of a candidate material relative to RHA
steel. For practical reasons, application of the Mass Equivalence Factor is
preferred,
CM
StSM FF
which directly indicates the relative mass reduction when replacing steel with a
lighter material for the same protection against the same projectile. Equating the
above terms yields (for the hydrodynamic approximation):
S
MF
F1
Figure 4 shows FS values versus the density ratio CM/ St. The candidate materials
are:
CM g/cm3
titanium 4.50
alumina 3.92
aluminum 2.80
glass fiber reinforced plastic (GFRP) 1.92
magnesium 1.75
polyvinyl chloride (PVC) 1.10
The round dots in Figure 4 represent the theoretical FM values for the specific
materials. In accordance with the hydrodynamic theory, PVC, as the lightest
material, yields the best mass equivalence protection factor.
66 Ceramic Armor Materials by Design
Figure 4. Space Equivalence Factor FS vs CM/ St
FIRST EXPERIENCES WITH AL2O3 CERAMICS
According to theory, the only ceramic material presented here is, obviously,
not very attractive. However, the experimental results obtained by firing at the
candidate materials with high precision shaped charges are more interesting, as
seen by the data notated by the triangles in Figure 4. Here the material strength
comes into play:
Due to the low strength of PVC, the measured value practically coincides with
the theoretical
Titanium reaches the same space equivalence, FS, as steel and a mass
equivalence FM = 1.744
GFRP has an astonishingly high FM = 3.3
Ceramics (alumina) need less space than steel with FS = 1.1 and at FM = 2.2, a
mass less than half the steel block
Ceramic Armor Materials by Design 67
This positive result becomes even better when using a reduced thickness. In
the described tests alumina blocks of 150-mm thickness were used. With blocks
of 40-mm thickness the mass equivalence mounts to FM = 4 and the space
equivalence to FS = 2. This proves that alumina is an ideal protection material
against shaped charges, being lightweight and needing little space. The reason for
this thickness-dependence is already described in 1 . It should be noted however
that the alumina blocks were housed in confinements.
The results of these investigations allow the conception of composite armor to
oppose shaped charges with values of FS 1 and FM 3 (Figure 5). Indeed, the
performance of such a target composition was confirmed in live 40-mm shaped
charge firing, the charge having a reference penetration of 180-mm RHA. For
correctness, the confinement mass of the side walls was neglected for the FM
calculation.
Next, the negative aspect of ceramic materials, namely their brittleness, had to
be dealt with. The area to be protected was subdivided into "tiles" in order to
confine the destruction of the ceramic armor to a small area. As described in 7,
8 two types of tile arrangements were tested against both model shaped charges
and heavy metal rods:
Triangular alumina tiles without damping elements at the boundaries and
The same tile shape with a knopped 1-mm thick rubber layer in between.
Figure 5. Composite armor: Aluminum
Confinement with Alumina/GFRP Inserts
68 Ceramic Armor Materials by Design
The damped version reduced the damage radius considerably (see Figs. 6, 7).
In the interim, this version was patented by RUAG Land Systems and model
armor specimens for trial firings were built 8 . In spite of these positive results,
no specific alumina applications were realized. On the one hand, this material is
rather expensive; on the other hand, there are other less delicate promising
materials which may protect against both shaped charges and KE projectiles.
Nevertheless, research studies continued on the protection effectiveness of
ceramics.
Figure 6. Damage to Ceramic Tiles
after Shaped Charge Impact with
Direct Contact between the Tiles
Figure 7. Damage to Ceramic Tiles
after Shaped Charge Impact with
Intermediate Layer between Tiles.
TEST RESULTS WITH SILICON NITRIDE (SI3N4)
The SI3N4 ceramic was provided by HTM of Biel/Bienne, Switzerland, under
the trademark "Hatemit". This firm offers this material as armor protection against
small arms projectiles 9 . Porous, hot isostatically pressed Si3N4 was tested
against large caliber shaped charges and KE projectiles 10 .
Against 105 mm KE rods, FM values between 2.5 and 3.2 were obtained,
nearly double the alumina values. However the extremely high price combined
with the reduced multihit capability have prevented practical applications so far.
Moreover porous Si3N4 with a density of only 2.3 g/cm3, requires a relatively high
armor thickness.
Ceramic Armor Materials by Design 69
TESTS WITH SILICON CARBIDE
In the Eighties, extreme protection requirements arose. One outstanding
example was the gunshield of a tank; the restricted depth required a maximum
space equivalence factor, at the same time the gun unbalance moment needed to
be kept to a minimum, resulting in a maximum mass equivalence factor
requirement. The only advantage of this optimization problem was the small
volume of the gunshield, which reduced teh probability of the multihit problem.
Cercom Inc. of Vista, California, USA offered their range of SiC materials
already in series production. This fact promised low prices, apart from the high FS
and FM values to be expected. The Cercom materials properties have been
described in 11 . SiC blocks of 150-mm x 150-mm x 30-mm were tested against
model shaped charges and KE rods.
Target 1 (Figure 8) consisted of two blocks in a 10-mm RHA confinement
with a 20-mm backing. Target 2 (Figure 9) was a combination of two SiC blocks
and 150-mm of GFRP in an identical confinement. RHA witness plates were
placed behind each target.
70 Ceramic Armor Materials by Design
Figure 8. Target 1
Figure 9. Target 2
CONCLUSIONS
The firing tests showed very high protection factors (see Table 1), the pure
SiC target naturally displaying higher values than the combined one.
Table 1: Test results (mean values)
L/D 20 Tungsten rod
Pref = 165-mm RHA
50- mm/50° Shaped charge
Pref = 330 mm RHA
Target 1: FM (SiC) 4.8 6.6
Target 2: FM (SiC + GFRP) 3.0 4.2
These good results encouraged us to try the 1:1 scale. Cercom provided SiC
blocks of 100-mm x 450-mm x 450-mm without problems. Targets of similar
layouts as the model targets were tested against large caliber rounds.
Concurrently, model tests were conducted at the Institute Saint Louis, France.
Partial results are published 12 . The protection capability of the full scale armor
was reduced (FM 3.2 instead of 4.8).
Table 1: Test results (mean values)
L/D 20 Tungsten rod
Pref = 165-mm RHA
50- mm/50° Shaped charge
Pref = 330 mm RHA
Target 1: FM (SiC) 4.8 6.6
Target 2: FM (SiC + GFRP) 3.0 4.2
Ceramic Armor Materials by Design 71
REFERENCES
1 H. G. Hopkins and H. Kolsky, "Mechanics of Hypervelocity Impact of
Solids", 4. Symposium on Hypervelocity Impact, 1960
2 G. Weihrauch, "Behaviour of copper rods impacting various materials with
velocities between 50 and 1650 m/s", ISL, Rapport - Bericht 7/71, 1971
3 C. L. Grabarek, "Penetration of Armor by Steel and High Density
Penetrators (U)", Ballistic Research laboratories, Aberdeen Proving Ground, Maryland, Memorandum Report No. 2134, October 1971
4 W. Odermatt, "Penetration Formula for Long Rod Penetrators", Defence
Procurement Agency, Report No. 1546, 14.04.2000
5 H.-J. Ernst and V. Wiesner and T. Wolf, "Armor Ceramics under High-
Velocity Impact of a Medium-Caliber Long-Rod Penetrator", Presented at Pac
RiM 4 Ceramic Conference in Maui/Hawaii, November 2001
6 R. Ochsenbein "Behaviour of Alumina targets impacted by shaped charge
jets", RUAG Munition (formerly Eidg. Munitionsfabrik Thun), Report No. X 010
042/1-67, 04.03.1980
7 R. Jeanquartier, "Behaviour of Alumina targets impacted by tungsten
rods", RUAG Munition (formerly Eidg. Munitionsfabrik Thun), Report No. FA X
010 027, 18.10.1979
8 R. Jeanquartier and B. Lehmann, "Firing tests with 35 mm APDS vs.
Composite targets", RUAG Munition (formerly Eidg. Munitionsfabrik Thun),
Report No. FA X 010 095, 08.04.1983
9 Dr. Hr. Thieme, "Silicon-Nitride as armor material against small calibre
munitions", RUAG Land Systems (formerly Eidg. Konstruktionswerkstätte Thun),
Report No. FB 00014, 14.02.1990
10 N. Schwizgebel, "Physical/Chemical Analysis of Silicon-Nitride",
Gruppe für Rüstungsdienst, Report No. FA-26-SIG Schw/ah-200/2270,
02.07.1984
11 Dr. H. Leber, "Material Properties of Silicon Carbide", RUAG Land Systems (formerly Schweiz. Unternehmung für Waffensysteme AG), Report No.
WTB 100009917, 11.09.2000
12 H.-J. Ernst and T. Wolf and W. Lanz, "SiC-Targets Against Differently
Scaled KE-Threats", RUAG Land, Report No. WTB 100009917, 11.09.2000
72 Ceramic Armor Materials by Design
STRUCTURE AND PROPERTIES OF SHOCK-RESISTANT CERAMICS
DEVELOPED AT THE INSTITUTE FOR PROBLEMS IN MATERIALS
SCIENCE, NAS OF UKRAINE
B.A. Galanov, O.N. Grigoriev, S.M. Ivanov and V.V. Kartuzov
Frantsevich Institute for Problems in Materials Science,
National Academy of Sciences of Ukraine
3 Krzhyzhanovsky St.
Kyiv, Ukraine 03142
ABSTRACT
The results of investigation of mechanical properties of a number of new
composite materials developed in IPMS, NAS of Ukraine are presented. A
prognosis of their ballistic properties was fulfilled on the base of the work [1].
INTRODUCTION
A wide range of ceramic materials, composites with ceramic matrix and
products made out of those has been developed practically for all fields of
economy (engineering, metallurgy, electrotechnics, chemical production,
environmental protection, etc.) and introduced by the Institute for Problems of
Materials Science, NAS of Ukraine within the period of 60-th of XX century and
up to present days. Actually, the Institute is one of those who developed bullet-
proof vests of all known protection classes. Bullet-proof vests production was
established by the Institute at a number of Ukrainian and Russian plants in the 80-
th.
R&D results on new ceramic materials, prospects for an employment in armor
and also for neighboring applications (wear-resistance and radiation protection)
are presented in this paper.
The materials under investigation are:
1. Ceramics and ceramic matrix composites (CMC) on the base of boron carbide;
2. Ceramics and CMC on the base of borides;
3. Ceramics and CMC on the base of silicon carbide;
4. Ceramic materials on the base of nitrides.
Ceramic Armor Materials by Design 73
To the extent authorized under the laws of the United States of America, all copyright interests in this publication are the propertyof The American Ceramic Society. Any duplication, reproduction, or republication of this publication or any part thereof, withoutthe express written consent of The American Ceramic Society or fee paid to the Copyright Clearance Center, is prohibited.
INVESTIGATIONS AND TECHNOLOGY
R&D was performed by:
Calculations and experimental determination of fields of internal stresses in
ceramics and CMC;
Optimization of structural and stress-strain states of CMC by an employment
of appropriate thermo-mechanical CMC model.
RAW MATERIALS
A wide range of refractory powder compounds supplied by different
manufacturers were used (Donetsk Plant of Chemical Reagents, Zaporozhye
Abrasive Combinat, Institute for Problems of Materials Science and H.C. Stark
(Germany)).
SINTERING
Within a framework of present R&D efforts a high-speed hot pressing
technology for products made out of different ceramics has been developed. The
hot pressing process was carried out on pilot installations with induction heating
in graphite molds without protective atmosphere.
CERAMICS PROPERTIES
Strength at Low Temperatures and Optimization of Structure and Composition of
Ceramics
Practically in all actual development programs, ceramics are not a single-
phase but represent some form of ceramic matrix composites. Elastic interaction
of phases at temperature and pressure changes during production and under
external thermo-mechanical effects results in a complex stress-strain state of a
material, which determines the features of its mechanical behavior. Therefore the
optimization of ceramics structure goes directly towards optimization of its fields
of internal stresses.
A mathematical formulation of fracture toughness criteria for the composite
with ceramic matrix and optimization methods for composite composition and
structure were proposed by B. Galanov and O. Grigoriev [2]. The introduction of
a high- component into the composite is accompanied by an increase of fracture
toughness. The maximum value of fracture toughness is shifted to the lower
second phase contents ( 10-30%) with an increase of CTE (coefficient of thermal
expansion) mismatch, the elastic characteristics, and grain sizes of the composite
phases. At high concentrations of the second phase, the minimum K1c with the
value of K1c 0 is found. It is caused by spontaneous failure under the effect of
thermal stresses with flaw size approximately equal to the grain size ( 10 m):
74 Ceramic Armor Materials by Design
For example, in SiC-TiB2 system, an improvement of composite properties
can be expected in the range of 20-30 vol. % of TiB2, and grain size of 5-10 m.
The analysis considers thermal stresses, partly generated by the difference in CTE
of the constituting phases of the composite.
The stress calculations present only an estimation of stress levels in phases,
without accounting for viscous-elastic relaxation. Therefore there is an important
role for experimental methods to determine internal stresses.
The experimental determination of thermal stresses gives us not only quantitative
information on internal stresses in phases, but also data on viscous-elastic
relaxation and the state of grain boundaries. A comparison of the experimental
and theoretical values of internal stresses shows a good agreement between theory
and experiment only if take a rather small T value ( T 1300 C) in the
calculation. It means that the temperature of the viscous-elastic transition is fairly
low (Tve 1300 C) and the composite possesses has a significant relaxation
ability at high temperatures (T 1300 C).
Boron Carbide Based Ceramics
The ceramic materials based on boron carbide are being employed thanks to
their high hardness, low volume mass, high effective capture cross-section of
neutrons and the like. Hot pressing is principal technique to obtain dense ceramic
out of boron carbide at temperatures 2000 – 2150 . That is why to produce the
materials on the base of boron carbide of special attention are methods of
activation of sintering. In present paper the activation of sintering process was
provided by borides additives (W2B5, TiB2 and ZrB2)
Composites of the B4C-TiB2 and B4C-ZrB2 Systems
Mechanical properties of the composites are considerably lifted up if the
second phase is of 15% (See Figure 1). The further lifting of borides
concentration brings to the reduction of fracture toughness and other
characteristics since the inner stress exceed an optimum level. Temperature
dependencies of hardness display that hardness of the composites on the base of
boron carbide at temperatures higher than 1100 exceeds the diamond and
boron nitride hardness (See Figure 2).
Composites of the B4C-W2B5 System
The ceramic was investigated in the volume content interval of tungsten
boride from 10 up to 90%, with a step of 10%. Table I displays the properties of
the materials presented.
Ceramic Armor Materials by Design 75
Table I. Properties of B4C-W2B5 system composites
Composition of
ceramics, vol. %
Density,
g/cm3
Bending
strength, MPa
HV, P=10N
GPa
90W2B5-10B4C 7.00 660 34
80W2B5-20B4C 6.75 500 35
50W2B5-50B4C 5.65 690 40
40W2B5-60B4C 5.10 590 38
10W2B5-90B4C 3.24 565 52
Figure 1. Variations of Young's modulus E, bending strength f, hardness HV,
fracture toughness 1 of ceramics on the B4C base versus second phase content:
– TiB2, – ZrB2
Figure 2. Temperature dependences of hardness for some superhard materials: 1-0
– B4C, 1-1 – B4C + 30% ZrB2, 1-2 – B4C + 40% ZrB2, 2 – diamond (Berkovich
indenter, = 70 ), 3 – diamond ( = 65 ), 4 – BN (hexanit), 5 – BN (elbor), 6 – BN
(PTNB)
The materials are obtained at comparatively low temperatures of hot pressing
(1800 C). Yet the mechanical properties, especially hardness, are turned to be
considerably higher than hot pressed boron carbide has (bending strength < 450
MPa, HV < 35 GPa).
Composites of the B4C-TiB2-W2B5 System
The materials of the above system apart from they have a high level of
mechanical properties (strength, hardness, impact- and wear-resistance) are
characterized by high linear coefficients of neutrons and -rays absorption.
Production conditions and properties of both one-phase materials – TiB2 and
W2B5, and composites TiB2-W2B5 and B4C-TiB2-W2B5 as well were studied.
Table II shows the properties of the materials obtained.
76 Ceramic Armor Materials by Design
Of particular interest is high strength of two- and three-phase materials (850-
1100 MPa) even at big grain size, when the synthesized powder was not grinded
before hot pressing.
Preliminary investigations have shown that the proposed material with its
ballistic and radiation protection characteristics exceeds the previously employed.
Table II.
Composition of
ceramics, vol. %
Density,
g/cm3
Bending
strength, MPa
HV, P=10N
GPa
TiB2 4,51 365 36
W2B5(W+B+Ni) 8.62 525 12
W2B5 8.37 590 24
50W2B5-50 TiB2 5.88 1110 26
70W2B5-25TiB2-5B4C,
Grain size-20 m
7.06 850 19.6
70W2B5-25 TiB2-5B4C,
Grain size-7 .m
6.76 900 20
Coarse Heterogeneous Composites Ceramic Metal
For operation under shock loading we developed materials with coarse
heterogeneous structure – granules from borides as wear resistant component,
binded with tough matrix.
The structure of materials is shown on the Figure 3, its strength is more than
500 MPa, hardness 15 GPa, density is in the range 4.8-4.9 g/cm3. Composites
advantages are related to ceramic carcass, providing resistance at high-speed
shock, while metallic component provides increased toughness of material and
products of them.
a) b)
Figure 3. TiB2 granules and coarse-heterogeneous composite after
abrasive-wear test
Ceramic Armor Materials by Design 77
DISCUSSION AND CONCLUSION
An analysis of penetration resistance was performed on the base of the
modified Alekseevskii-Tate model for nonstationary penetration of long rods into
targets [1].
“Static” component of resistance to penetration into target Rt was evaluated.
Total resistance to penetration is defined by three factors: static Rt, kinematic Pk
(0.5 U2) and dynamic Pd, relative contributions of those are varying during the
penetrator – target interaction.
In accordance with the accepted model the value Rt is defined by the system of
elastic and strength characteristics of target material: , E, , Y, f, etc. Table III
shows some of those.
0 1 2 3 4 5 6
-20
0
20
40
0.0 0.1
-20
0
20
Rt
Pk
Pd
Pc= Rt+Pk+Pd
Penetration, mm
Pressure, GPa
Figure 4. Structure of penetration resistance (contact pressure Pc) in Al2O3.
Solid line – data from [1], dashed – from table III. Impact velocity 1600 m/s (steel
penetrator).
78 Ceramic Armor Materials by Design
In case target material possesses a tangible plasticity value Y that defines
material behavior on the boundary plastic — elastic
material is the yield stress of this material. In the
case of brittle material, the comminuted (pulverized)
zone is formed in the contact area as a result of
multiple fragmentation and value Y is defined by the
strength characteristics of material and is obviously
close to strength limit under compression. It was
established that at least for porous ceramic H
compression. Analogously, during quasi-static
indentation into brittle materials, semi-spherical
fragmentation zone is formed in the area of
deformation core (See Figure 5) with a pressure on
its boundary = Y, that defines the impression size
and, consequently, contact pressure and material's
hardness.
Figure 5. Core structure in
TiN/AlN ceramics
Ceramic Armor Materials by Design 79
Table III. The properties of developed ceramic materials Materials Additions,
second
phases,
vol.%
Density,
g/cm3
Young's
modulus,
E, GPa
Hardness,
HV
(P=5N),
GPa
Yield
stress,
Y, GPa
E/Y Strength, fbend,
RT,MPA
Penetration
hardness
HP, GPa
HP B4C - 2.5 450 30 20 22.5 500 2.65
HP
B4C/MeB2
ZrB2 or
TiB2(up to
30%)
up to 2.7 460 35 23 20 500-800 2.66
HP
B4C/CaB6
CaB6 up to
100%
2.5 450 25-30 17-20 25 400-500 2.65
HP CaB6 - 2.4 460 25 17 26 400 2.63
TiB2 - 4.5 550 27 18 30 500-700 2.68
TiB2/CaB6 CaB6 up to
30%
4.4 530 45 30 17 700 2.68
W2B5
W2B5
(1%Ni)
- 10-13
10-13
775
775
30
12
18
8
43
97
500
500
3.05
3.73
TiB2/
W2B5/
B4C
- 4-10 600 30 20 30 800-1000 2.68
3.0
TiB2/
W2B5
W2B5 up to
50%
4.5-10 600 35 23 26 700-1000 2.68
2.86
HP SiC B4C up to
5%
3.2 460 20 13 35 300-500 2.72
HP SiC/
MeB2
ZrB2 or TiB2
(up to 30%)
3.3 450 25 17 26 600-700 2.66
RS
SiC/MeB2
The same 3.2 440 20 13 34 500 2.7
HP
TiN/AlN
AlN up to
100%
3.2 400 20 13 31 400-500 2.7
HP AlN - 3.2 280 12 8 35 300-400 3.07
HPSi3N4/
ZrO2/Y2O3
ZrO2 (up to
30%), Y2O3
(up to 10%)
3.3 320 16 12 26 700 2.78
S -
SIALON
z = 2-5 3.15 220 14 9 24 500 2.93
HP -
SIALON
AlN/Y2O3
(up to10%)
3.2 350 21 14 25 600 2.73
Al2O3 - 4 400 15 10 40 500 3.04
Al2O3 [1] - 3.5 373 2.62 142 262 7.45
Al2O3/
TiB2/
ZrO2
TiB2 up to
30%, ZrO2
up to 30%
4.5 430 16 12 35 800 2.96
80 Ceramic Armor Materials by Design
In correspondence with indentation models of Tanaka, when radius of semi-
spherical core approximately equals to the radius of contact area, the ratio of
hardness to Y is ~ 1, which allows to evaluate the upper boundary of Y with
hardness value. The lower boundary of Y is defined evidently by the relationship
HV~3Y for plastic material.
The authors are also of the opinion that if under impact radial cracks are
formed outside the comminuted zone then their formation must be defined not by
material's tensile strength but by «contact strength» of material tested precisely
under the contact loading and which can be evaluated using the value of length of
radial cracks around the hardness indents. The technique of contact strength
measurement is presented in [3].
Resistance to penetration was characterized by the penetration work per the
unit volume of extruded material: "penetration hardness" HP= 1/P PcdP (P –
depth of penetration).
The analysis of the results has shown that under investigated conditions of
impact the following conclusions can be made:
1. During penetration the pressure on the contact surface increases.
2. High-strength ceramics in the wide range of variation of its characteristics
demonstrates small change in the "penetration hardness" – 2.5–3 GPa.
3. "Penetration hardness" increases with the growth of parameter E/Y (up to 7
GPa at E/Y=140), however, the depth of penetration also sharply increases.
ACKNOWLEDGMENT.
The authors would like to acknowledge the support from ARL under the
contract 68171-01-M-5848 and scientific coordinator Dr. W. Gooch.
REFERENCES
1. B.A. Galanov, S.M. Ivanov, and V.V. Kartuzov. "On one new
modification of Alekseevskii-Tate model for nonstationary penetration of long
rods into targets". Proc. of HVIS'2000, Journal of Impact Engineering, 26 201-10
(2001) (to be published).
2. B.A. Galanov, O.N. Grigoriev, and V.I. Trefilov, "Ceramic Matrix
Composites: theoretical fundamentals", in Ceramic- and Carbon-matrix
Composites, Edited by V.I. Trefilov, Chapman & Hall, 3-29, 1995.
3. B.A. Galanov, O.N. Grigoriev, and E.G. Trunova, "Contact strength and
statistikal fracture mechanics of ceramics"; p.19 in Proc. of Int. Conf. “Current
Problems of Strength”, 3-5 July, 2001, Kiev.
Ceramic Armor Materials by Design 81
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CERAMIC ARMOR WITH SUBMICRON ALUMINA AGAINST
ARMOR PIERCING PROJECTILES
E. Strassburger, B. Lexow
Fraunhofer-Institut für
Kurzzeitdynamik Ernst-Mach-Institut
(EMI)
Am Klingelberg 1
D-79588 Efringen-Kirchen, Germany
A. Krell
Fraunhofer-Institut für Keramische
Technologien und Sinterwerkstoffe
IKTS, Winterbergstr. 28
D-01277 Dresden, Germany
ABSTRACT
In a joint project of the Fraunhofer Institute for Ceramic Technologies and
Sintered Materials (IKTS) and the Ernst-Mach-Institute (EMI), aluminum oxide
ceramics with submicron grain size were developed and tested ballistically. In
DOP-tests with tungsten alloy projectiles, the new ceramics revealed a ballistic
efficiency superior to commercial alumina grades.
Additionally, the ballistic performance of the new submicron and commercial
alumina against armor piercing (AP) steel core projectiles was investigated. The
ceramic/aluminum targets were also tested in a Depth of Penetration (DOP)
configuration. The influence of ceramic layer thickness and sequence was
determined with laminated targets.
INTRODUCTION
High strength ceramics are employed as ballistic protection material when a
high protective strength is required at a low weight. In order to improve the bal-
listic performance of a ceramic, it is necessary to know the correlations between
the microstructure and the ballistic resistance. However, for studying the correla-
tions between microstructure and ballistic resistance of one type of ceramic, it is
essential to have well defined materials, where individual parameters like grain
size, purity, porosity and density can be adjusted with high accuracy. The results
of previous studies1 have indicated, that only high purity ceramics with relative
densities > 98.5 % should be used in investigations on the influence of grain size
and hardness on ballistic performance. On one hand, earlier tests have demon-
strated that the ballistic resistance of ceramics increases with increasing
hardness2. On the other hand, it is known that the hardness of polycrystalline
Ceramic Armor Materials by Design 83
To the extent authorized under the laws of the United States of America, all copyright interests in this publication are the propertyof The American Ceramic Society. Any duplication, reproduction, or republication of this publication or any part thereof, withoutthe express written consent of The American Ceramic Society or fee paid to the Copyright Clearance Center, is prohibited.
ceramics increases with decreasing grain size3. Therefore, high purity aluminum
oxide ceramics with sub-µm grain size were developed and ballistic tests were
conducted in a collaboration between IKTS and EMI. The first part of the
investigations comprised materials with 1 µm, 0.5 µm and 0.3 µm grain size.
The material with 1 µm grain size was supplied by Dornier GmbH, Germany. The
sub-µm grain size materials were developed by IKTS. The commercially
available AD995 (trade name CAP-3) of Coors, Golden, Colorado, was used as
reference material because it had exhibited the highest ballistic mass efficiency
among previously tested commercial alumina grades. The ballistic resistance of
these materials was tested in a Depth of Penetration (DOP)-configuration with an
armor steel backing by means of tungsten alloy projectiles.
A second part of the study focuses, on one hand, on the investigation of sub-
µm alumina with improved strength, due to improved manufacturing processes
which lead to a significant reduction of the number of microscopic flaws. With
respect to light armor applications of the materials, the objective of the
investigations was to determine the potential of the materials for the defeat of
steel core projectiles and possible armor efficiency improvements by laminated
target configurations.
MATERIALS
The relative density, grain size, hardness and bending strength of the tested
materials are specified in Table I. The materials designated “S-“ were manufac-
tured at IKTS by means of spray drying, cold isostatic pressing and unpressurized
sintering in air. The “D-0.9” material was supplied by Dornier GmbH, Frie-
drichshafen, Germany. A more detailed description of the materials is provided in
an additional paper by A. Krell4.
Table I. Material specifications
Relative density Grain size Hardness
HV10
4-point
bending strength
(%) (µm) (GPa) (MPa)
S-0.3 92.5 0.32 15.0 Not determined
S-0.5 99.3 0.53 19.3 203 16
S-0.7 99.5 0.71 19.1 526 55
D-0.9 98.7 0.92 15.7 244 41
AD995 (CAP-3) 98.8 10-20 12.3 350 25
BALLISTIC TESTING
Tungsten projectiles versus ceramic/steel targets
In the first part of the investigations, different types of alumina were tested in
a DOP-configuration with a RHA (Rolled Homogeneous Armor steel) backing of
84 Ceramic Armor Materials by Design
hardness 300 HV30. A tungsten alloy cylinder with a hemispherical nose was
used as projectile. The diameter of the projectile was 10 mm, the length 32 mm
and the mass was 44 gram. The impact velocity was 1250 m/s nominally. As a
figure-of-merit for ballistic performance, the ballistic mass efficiency, Em,, was
chosen, which is determined from the residual penetration PR, the penetration into
the reference steel target, Pref, the thickness of the ceramic, TCer and the densities
St, Cer of the steel and the ceramic. Figure I shows a schematic of the test
configuration. The definition of the mass efficiency is given in equation (1).
Figure I. Schematic of DOP-Test configuration
RStCerCer
refStm
PT
PE (1)
The DOP data of all tested materials are presented in Figures II and III. On the
left hand side, the residual penetration is shown as a function of ceramic
thickness, whereas on the right hand side, the mass efficiency is plotted versus the
ceramic weight fraction Fcer/ Ftot = cerTCer/( cerTCer + StPR). The diagram for the
reference material AD995 (Fig. II) exemplifies the behavior observed with all
commercial aluminas. A linear decrease of residual penetration as the ceramic
thickness increases is associated with a linear increase of mass efficiency with
increasing ceramic weight fraction. A linear extrapolation to the point, where the
projectile is stopped just at the ceramic-steel interface (PR = 0), yields the
maximum mass efficiency Em,max for that material. For AD995 Em,max was 2.1
with the projectile/ target combination considered here. The sub-µm alumina
exhibited a different behavior (see Fig. III). Compared to AD995, a slightly
higher residual penetration PR was observed at a ceramic thickness up to 10 mm.
However, with a ceramic thickness of more than 15 mm, the residual penetration
decreased rapidly, which implies a higher ballistic efficiency. For this sub-µm
Ceramic Armor Materials by Design 85
material, both dependencies, PR-TCer and Em versus ceramic weight fraction can
be approximated by a second order polynomial fit.
Figure II. DOP data with AD995 (CAP-3)
0
10
20
30
0 10 20
TCer [mm]
PR [
mm
]
30
AD995, monolithic
AD995, 6.7 mm + 15 mm
AD995, 15 mm + 6.7 mm
1,0
1,4
1,8
2,2
2,6
3,0
0,0 0,2 0,4 0,6 0,8 1,0
FCer/ Ftot
Em
AD995, monolithic
AD995, 6.7 mm + 15 mm
AD995, 15 mm + 6.7 mm
Figure III. DOP data with sub-µm aluminas
0
10
20
30
0 10 20
TCer [mm]
PR [
mm
]
30
D-1.0
S-0.7
S-0.5
S-0.3
Fit S-0.5
black symbols: monolithic and
10 mm + 10 mm
grey symbols: 5 mm + 15 mm
open symbols: 15 mm + 5 mm
1,0
1,4
1,8
2,2
2,6
3,0
0,0 0,2 0,4 0,6 0,8 1,0
FCer/ Ftot
Em
D-1.0
S-0.7
S-0.5
S-0.3
black symbols: monolithic and
10 mm + 10 mm
grey symbols: 5 mm + 15 mm
open symbols: 15 mm + 5 mm
The extrapolation of the Em-curves for D-0.9, S-0.7 and S-0.5 resulted in
significantly higher Em, max values compared to AD995. Extrapolation, based on
the results with 10 mm + 10 mm targets, yielded a maximum mass efficiency of
2.6, and even 2.9 could be achieved with the 5 mm + 15 mm configuration for
each of the three materials.
Two important conclusions can be drawn from the presented data. The results
with D-0.9, S-0.7 and S-0.5 show that hardness is much more important for the
86 Ceramic Armor Materials by Design
ballistic efficiency than bending strength. The fact that Em,max of the very fine
grained S-0.3 material is higher than with AD995, but significantly lower
compared to the coarser, but dense, ceramics S-0.5 and S-0.7 demonstrates, that
there is no separate influence of grain size on Em,max beyond the hardness effect.
An additional effect was observed with laminated ceramic targets of 20 mm
total thickness. The lowest residual penetration was observed with the 5 mm + 15
mm targets, the highest PR occurred with the 15 mm + 5 mm plate sequence. The
lamination effects can be explained qualitatively as follows. When the projectile
first hits a thin front layer, this layer will be fragmented very rapidly and will
exhibit only a low ballistic resistance. However, pre-damage to the second plate
will be reduced so that the projectile has to penetrate a material with higher
ballistic resistance compared to the case of a monolithic target. When the thick
plate is at the front, the penetration of this plate is like that in a monolithic target.
However, the thin plate at the back will be shattered by several wave reflections
that results in a strongly reduced ballistic resistance of this plate. With two layers
of equal thickness both effects, stronger pre-damage of the first plate and reduced
pre-damage of the second plate, appear to compensate one another.
Steel core projectiles versus ceramic/aluminum targets
Three types of alumina were tested in a DOP-configuration with armor
piercing (AP) steel core projectiles of 14.5 mm caliber at an impact velocity of
1045 15 m/s. The projectiles had a total mass of 64.1 g, whereas the mass of the
steel core was 40.5 g. Aluminum (AlCuMg1) of tensile strength 400 MPa was
used as backing material. The ceramic tiles were glued to the backing by means of
the polyurethane glue Sikaflex and the joint between the lateral steel confinement
and the edge of the ceramic was filled with epoxy.
Figure IV shows the residual penetration versus ceramic thickness, not only
for monolithic targets of AD995, D-0.9 and S-0.7, but also for laminated targets
consisting of two or three plates of equal thickness or two plates with a thickness
ratio of ½. The dashed line indicates the overall behavior with monolithic ceramic
targets. Two sections can be distinguished where the decrease of PR with
increasing ceramic thickness can be approximated linearly. However, for ceramic
thickness 10 mm < TCer 15 mm, the slope is only half of that in the range from 5
mm to 10 mm thickness. When TCer is 5 mm or less, the ballistic resistance is very
small as the results with D-0.9 and S-0.7 indicate. In the range from 5 mm < TCer
15 mm, no significant difference was observed with the three types of alumina.
That means, with respect to the ballistic resistance, no benefit of the increased
hardness and strength compared to AD995 was achieved for this type of ballistic
test with D-0.9 and the high strength sub-µm grain size S-0.7. This result
indicates, that there is a saturation with respect to the influence of hardness on the
ballistic resistance against steel core projectiles. When the hardness is sufficiently
Ceramic Armor Materials by Design 87
high erosion and/or break up of the steel core is initiated. A higher hardness does
not lead to a more efficient erosion of the steel core. The fact, that there was no
benefit of the improved strength could be attributed to the design of the DOP test
configuration. The quasi semi-infinite backing prevents or reduces bending of the
ceramic specimen during the first phase of the projectile/target interaction. The
role of the strength of the undamaged material decreases during penetration as the
pre-damage of the material by stress waves progresses. However, that does not
exclude benefits of the high strength in “thin” targets, where bending of the
backing occurs.
Figure IV. Residual penetration versus ceramic thickness
0 5 10 15 20
TCer [mm]
AD995, monolithic
AD995, 2 x 6.7 mm, plates loosely stacked
AD995, 2 x 6.7 mm, plates glued with Sikaflex
AD995, 2 x 6.7 mm, plates ground on both sides, stacked
D-0.9, monolithic
D-0.9, 3 x 5 mm, glued with Sikaflex
D-0.9, 5 + 10 mm, glued with Sikaflex
S-0.7, monolithic
arb
itra
ry u
nit
s
With AD995, three different target types consisting of two plates of 6.7 mm
thickness were assembled and tested. The targets differed in the flatness of the
ceramic plates and in the way the components were joined. All two layer targets
exhibited a significantly lower ballistic resistance compared to the resistance
expected for monolithic targets of the same total thickness. The highest residual
penetration was observed when the ceramic plates were only loosely stacked.
Joining the two ceramic plates with polyurethane glue resulted in a reduction of
PR. However, the best performance of the two layer targets was observed when the
two plates were ground on both sides and stacked without glue. This observation
implies that the ballistic resistance against AP projectiles is the better the closer
the target is to the monolithic case.
Targets consisting of three layers of equal thickness (3 x 5 mm) and two layer
targets (5 mm + 10 mm) of D-0.9 were tested. In both configurations, the ceramic
88 Ceramic Armor Materials by Design
plates were glued together with Sikaflex and the bottom ceramic plate was glued
to the aluminum backing. The 5 mm + 10 mm targets performed significantly
worse than the monolithic targets, whereas the ballistic resistance of the three
layer targets was heavily degraded compared to the monolithic ones. The relative
performance of the laminated and monolithic targets is summarized in Figure V,
where it is presented as the ratio of the residual penetrations with monolithic and
laminated targets (PR, monolith/PR, laminate). Since the ballistic resistance of single 5
mm plates of D-0.9 was very low, the poor performance of the 3 x 5 mm targets
could be expected.
Figure V. Relative performance of laminated and monolithic targets
0
0,2
0,4
0,6
0,8
1
stacked glued ground, no glue monolithic
AD995
2 x 6.7 mm
0
0,2
0,4
0,6
0,8
1
3 x 5 mm 5mm+10mm monolithic
D-0.9
The efficiency of the ceramic is connected to its ability to erode and break up
the penetrating steel core of the projectile. Thus, the residual mass mR of the steel
core is also a measure of the efficiency of the armor. In Figure VI, the residual
mass of the steel core is plotted versus ceramic thickness. With the monolithic
targets, a strong decrease of mR was observed at a ceramic thickness of 10 mm.
At lower ceramic thickness, typically, an eroded steel core of mR > 30 g was left,
whereas with TCer > 10 mm, only 5-10 fragments of the steel core with a total
mass of less than 15 g could be found (see Fig. VI).
SUMMARY
The influence of grain size, hardness and strength on the ballistic performance
of Al2O3-ceramics was determined by means of DOP-tests.
Tungsten alloy projectiles were employed in order to determine the ballistic
resistance of ceramic/steel targets.
In this projectile/target combination, alumina with sub-µm grain size exhibited
significantly higher maximum mass efficiencies than commercially available
alumina tested under the same conditions.
The results clearly indicate that there is no separate influence of grain size or
of flaws beyond their impact on hardness.
The ballistic resistance against tungsten projectiles can be increased by
laminated targets consisting of a thin front layer and a thick rear layer.
Ceramic Armor Materials by Design 89
14.5 mm AP steel core projectiles were used in order to assess the ballistic
resistance of ceramic/aluminum targets. No significant difference was observed in
the penetration behavior of the three alumina types tested.
That means, with respect to the ballistic resistance no benefit of the increased
hardness and strength compared to AD995 was achieved with D-0.9 and the high
strength sub-µm grain size S-0.7.
Lamination of the ceramic targets resulted in significant losses of ballistic
performance.
Figure VI. Residual steel core mass and state of the residual projectiles
0
10
20
30
40
50
0 5 10 15 20
TCer [m m ]
mR [g
AD995, monolithic
AD995, 2 x 6.7 mm, stacked
AD995, 2 x 6.7 mm, glued
AD995, 2 x 6.7 mm, ground,no glue
D-0.9, monolithic
D-0.9, 3 x 5 mm
D-0.9, 5 mm + 10 mm
S-0.7, monolithic
TCer, monolithic
< 10 mm
TCer, monolithic
> 10 mm
REFERENCES1B. James, “The influence of the material properties of alumina on ballistic
performance,” pp. 3-9 in Proceedings of the 15th
International Symposium on
Ballistics (Jerusalem/Israel, 1995 published by the Organizing Committee).2I. Faber, K. Seifert and L.W. Meyer, “Correlation between the mechanical
data of ceramics and their protective power against impact loading” (in German),
Final Report EB 6/95 (part 3), Technical University Chemnitz-Zwickau,
Department of Engineering Materials, 1995. 3A. Krell and P. Blank, “Grain Size Dependence of Hardness in Dense
Submicrometer Alumina,“ J. Am. Ceram. Soc. 78 4 1118-20 (1995).4A. Krell and E. Strassburger, “High-Purity Submicron -Al2O3 Armor
Ceramics – Design, Manufacture, and Ballistic Performance”, Proceedings of
PAC RIM IV, Ceramic Armor Materials by Design (Wailea, Maui, Hawaii,
2001).
Figure VI. Residual steel core mass and state of the residual projectiles
0
10
20
30
40
50
0 5 10 15 20
TCer [m m ]
mR [g
AD995, monolithic
AD995, 2 x 6.7 mm, stacked
AD995, 2 x 6.7 mm, glued
AD995, 2 x 6.7 mm, ground,no glue
D-0.9, monolithic
D-0.9, 3 x 5 mm
D-0.9, 5 mm + 10 mm
S-0.7, monolithic
TCer, monolithic
< 10 mm
TCer, monolithic
> 10 mm
90 Ceramic Armor Materials by Design
ARMOR ALUMINA CERAMICS
Eugene Medvedovski
Ceramic Protection Corporation
3905 – 32nd
Street N.E.
Calgary, Alberta, T1Y 7C1, Canada
ABSTRACT
Dense alumina ceramics are still one of the most cost-effective armor
materials among different structural ceramics used for ballistic protection. They
have high mechanical properties and excellent manufacturability. High-alumina
ceramics with an Al2O3 content ranging from 97 to 99.6-wt.% and alumina-
zirconia ceramics produced by Ceramic Protection Corporation have been
successfully used as an armor material for personnel and vehicular ballistic
protection. They are manufactured by slip casting and pressing technologies
depending on the required shape and quantity. The main properties of the
ceramics, which affect ballistic performance, and ballistic test results are
examined and analyzed. Only the combination of all physical properties and
microstructure, as well as the optimization of the manufacturing process, should
be considered for selection and evaluation of armor ceramics.
INTRODUCTION
Ceramic armor was originally developed for “bulletproof vests” and seat-
armor in helicopters. At the present time ceramic armor is mainly used for
personnel and vehicular ballistic protection in military forces and by tactical
police teams, for protection of some critical parts of aircraft and helicopters and
for blast protection against landmines.
The mechanisms of ballistic protection for ceramic and metal armor are
significantly different. Metals absorb the energy of projectile by a plastic
deformation mechanism. In the case of ceramics, the kinetic energy of the
projectile is absorbed by a fracture energy mechanism. Usually the ceramic armor
system consists of the monolithic ceramic or composite ceramic-metal body
covered by ballistic nylon and bonded with a high tensile strength fiber lining
such as KevlarTM
, SpectraTM
or fiberglass. Also some soft metals (e.g. aluminum
thin sheets) may be used as a backing material. Upon impact of the bullet
Ceramic Armor Materials by Design 91
To the extent authorized under the laws of the United States of America, all copyright interests in this publication are the propertyof The American Ceramic Society. Any duplication, reproduction, or republication of this publication or any part thereof, withoutthe express written consent of The American Ceramic Society or fee paid to the Copyright Clearance Center, is prohibited.
(velocity greater than 700-800 m/sec), the hard-facing ceramic body is cracked
and broken, and the residual energy is absorbed by the soft reinforced backing
material. This backing material also must support post-impact fracturing of the
ceramic body caused by the bullet and the bullet itself.
Consideration of ballistic protection systems must take into account several
factors: the type of ballistic threat, the ability to manufacture the armor system
and the properties of the armor system components. These include such factors as
threat level, multi-hit performance, environmental conditions, space limitations,
manufacturing challenges, cost and weight limitations, physical properties of both
facing and backing material and the overall ballistic performance of the system.
Different armor ceramics including monolithic ceramics and ceramic-matrix
composites are described in [1-7]. Among them alumina ceramics are of low cost
and may be manufactured using a variety of methods, i.e. slip casting, pressing,
injection molding and some others, without the use of expensive equipment, e.g. a
kiln with special controlled atmospheres. Despite elevated density (up to 3.95
g/cm3), alumina ceramics are used for ballistic protection. In this paper, the high-
alumina and alumina-zirconia ceramics commercially produced by Ceramic
Protection Corporation (CPC) are reviewed and studied. Different armor products
(tiles and monolithic curved plates) are manufactured from these ceramics with a
high quantity (e.g. several hundreds plates and several thousand tiles per day).
They are designed and manufactured in accordance with the specific customer
demands depending on the application, required performance and level of
protection; they may be obtained as bare ceramic products, or they may be laid-up
with backing materials. They have been successfully used for personnel and
vehicular ballistic protection.
EXPERIMENTAL
Materials
The studied armor alumina ceramics are based on the systems Al2O3-SiO2-
CaO-MgO and Al2O3-MgO with an Al2O3 content approximately 97, 98, 98.5 and
99.6-wt.%. The alumina-zirconia ceramics is based on a specially selected ratio
between Al2O3 and ZrO2 (Y2O3 is used as a stabilizing agent). The starting
alumina powders producing by Pechiney - Altech (France) and Alcoa World
Chemicals (USA) have a high purity (minimum 99.8-wt.% of Al2O3); the -form
content is 95-wt.% or greater depending on the alumina grade. The average
median particle size and crystal size of the aluminas range from 0.35-0.45 to 1.1-
1.4 µm, and their specific surface BET ranges from 8-11 to 2.8-3.3 m2/g for the
used alumina powders, respectively. In the case of the alumina-zirconia ceramics,
the zirconia powder producing by Tosoh Corp. (Japan) is also used as a raw
material. The partially Y2O3-stabilized zirconia powder has the median particle
and crystal size of 0.3-0.4 µm and the specific surface BET of 8-9 m2/g.
92 Ceramic Armor Materials by Design
Manufacturing
Manufacturing methods used for production include slip casting and dry
pressing processes depending on the shapes and quantity of ceramic products to
be made. Experimental, pilot-scale and production studies allowed for
optimization of the following manufacturing steps:
Ceramic water-based slip preparation depending on the batch composition
(consisting up to 77-80-wt.% of solid), including the development of the
dispersant and binder systems;
Slip casting process providing manufacturing of single, double, and triple
curve plates with the custom designed shape and dimension;
Spray drying process providing a powder yield of up to 96%, press-powder
preparation providing “donut”-free spherical particles with adjustable sizes;
Uniaxial pressing process;
Drying and firing processes (firing temperature less than 1550oC), including
kiln loading, depending on the shape and size of the products and the optimal
firing curve;
Bonding process of ceramics with a backing material including adhesive
preparation, Kevlar, fiberglass and nylon preparation, thermal treatment of the
glued ceramic product with a backing material in an autoclave where
temperature, pressure, and vacuum are applied;
Quality control system which provides overall quality control and possible
adjustments at each manufacturing step.
The manufacturing of armour products at CPC is ISO 9002 certified. Each
step of the manufacturing process must be accompanied by the corresponding
quality control procedures. The quality control starts with the raw materials
verification. The following ceramic manufacturing parameters are controlled;
some are adjusted individually in order to achieve required parameters:
Sequence and duration of the starting materials mixing and milling, specific
gravity, viscosity, pH of the initial slips;
Binder and plasticizer component contents, sequence of addition to the initial
slip, specific gravity, viscosity, pH of the resultant slips (particularly if the slip
is used for slip casting or for spray drying), and casting rate of the slip if a
new lot of raw materials is started using for manufacturing;
Spray drying parameters (air pressure, inlet and outlet temperatures, flow rate,
etc.);
Granulation process parameters;
Particle size distribution, bulk density, powder flow rate, moisture content for
spray dryed powders and for powders ready-to-press (RTP); compression
coefficient for the RTP-powders;
Ceramic Armor Materials by Design 93
Slip casting and pressing processes and parameters;
Dimensions and weight for green products;
Firing parameters, including firing curve and final firing temperature, oxygen
level, air pressure, and kiln loading.
The following parameters are tested for the fired ceramics:
Fired and total shrinkage, dimensions and shape parameters (curvature for the
body armor and special vehicular armour plates, flatness for tiles);
Density and open porosity for the products and the witness samples made
from the same material as products from each kiln fired at different levels of
the kiln;
Physical properties (Vickers hardness, fracture toughness, sonic velocity,
Young’s Modulus, flexural strength) for the witness samples from the
selected firings using the specially developed testing protocol;
Ballistic performance for the products in accordance with testing protocols.
Testing
Microstructure was studied using transmission and scanning electron
microscopes. Density, porosity, and water absorption were tested using the water
immersion method based on Archimedes law. Four-point flexural strength was
tested in accordance with ASTM C1161. Young’s Modulus and sonic velocity
were tested by the ultrasonic technique measuring the longitudinal ultrasonic
velocity in accordance with ASTM C769 and by the resonant frequency method
in accordance with ASTM C885. Vickers hardness was tested in accordance with
ASTM C1327 at loads from 0.3 to 50 kg; the load 10 kg was used as a “standard”.
Fracture toughness KIc was also determined using the indentation technique under
the load of 10 kg. The test samples with required dimensions were cut from actual
products or from the test tiles produced by the mentioned technologies.
Ballistic performance of ceramics bonded with appropriate backing materials
was tested in accordance with the NIJ 0101.03 and NIJ 0101.04 standards using
the weapons M16, KAR 98K, AK47 and some others (caliber 0.30). Depending
on the application and the required level of protection, the ammunition 7.62x51-
mm NATO Ball Full Metal Jacket (FMJ) with a lead core, 7.62x63-mm Armor
Piercing M2 FMJ with a tungsten carbide core, 7.62x39-mm Russian Ball FMJ
with a steel core and some others were used. Depending on the ammunition, the
bullet weight, velocity and energy are varied. The bullet velocity was controlled
using a chronograph. The trauma after shooting was measured using a Roma
Plastilina modeling clay supported armor system on the back; the trauma in clay
duplicated the trauma in armor. The damage zone of the ceramics, including
ceramic fragmentation and the bullet were observed. Considering ceramic armor
systems for ballistic testing, the flat tiles (100x100 mm or greater) with a
thickness of 7-15 mm were used for the single shot testing. Also, the tiles (50x50
94 Ceramic Armor Materials by Design
or 100x100 mm) assembled as a flat panel, and flat tiles (155x200x8-9 mm), as
well as the actual plates with different configuration, were used for multi-hit
ballistic testing (with approximately 50 mm spacing between hits).
RESULTS AND DISCUSSION
Microstructure and Physical Properties
All studied ceramics are fully dense (water absorption is not greater than
0.02%) after firing at temperature less than 1550oC. Phase composition and
microstructure of the AL97ML, AL98 and AL98.5 alumina ceramics (a number in
the ceramic composition denotes an approximate Al2O3 content) are similar, and
they consist of corundum grains (the major phase) bonded by a small amount of
anorthite crystals and a silicate-based glassy phase. A small amount of mullite
crystals is also present in the AL97ML ceramics. The AL99.6 alumina ceramics
consists of corundum grains bonded by spinel crystals and a very small amount of
a glassy phase that formed due to the presence of oxides-impurities.
The ultimate grain size of the alumina ceramics depends on the initial batch
composition, initial particle size and particle size distribution of the starting
alumina powders. As expected, as an alumina powder with a smaller particle and
median crystal size was used, a fine-crystalline structure with a smaller grain size
was achieved. The average corundum grains are ranged from 1-3 µm for the
AL99.6 (mostly isometric) to 3-6 µm (isometric) and (2-3)x(5-8) µm (short
prismatic) for the AL97ML ceramics. A glassy phase is distributed between
grains uniformly and, as expected, the amount of a glassy phase increases as the
alumina content decreases.
The alumina-zirconia AZ ceramics based on the special ratio between alumina
and partially stabilized zirconia (PSZ) does not have a glassy phase; zirconia
grains with a size less than 1 µm are uniformly distributed between corundum
grains with a size of 1-2 µm. The zirconia phase, probably, inhibits the corundum
grain growth during sintering.
All these microstructure features affect physical properties and ballistic
performance of the ceramics. Physical properties depend on the Al2O3 content, the
size and shape of corundum grains, the amount, composition and distribution of a
glassy phase cemented the crystalline phase, the presence and composition of the
“secondary” crystalline phases, and closed porosity. They also depend on the
“stressed conditions” at the boundary of the corundum grains and a glassy phase.
These factors are governed by the wetting of alumina particles by a liquid phase
and by the interaction between them during sintering, firing and cooling, as well
as by the difference in thermal expansion between crystalline and glassy phases.
The major properties of the studied ceramics are presented in Table 1.
Young’s Modulus, sonic velocity, and flexural strength of the studied alumina
ceramics tend to increase as the Al2O3 content increases and with a smaller grain
Ceramic Armor Materials by Design 95
size. For example, the notable difference between Young’s Modulus and sonic
velocity data for the AL98 and AL98.5 ceramics, despite their closeness in Al2O3
content, can be explained by a smaller grain size, a higher densification, a higher
uniformity of microstructure and lower closed porosity for the AL98.5 ceramics.
The AZ ceramics demonstrates the highest value of flexural strength (>500 MPa)
due to the presence of the PSZ phase and fine-crystalline structure.
Hardness depends on the composition and microstructural features and also on
the load used for measuring. As the load is higher, as HV number is lower.
Hardness values tend to increase as the Al2O3 content and the corundum grains
content increase for the studied alumina ceramics. In the case of a higher glassy
phase content, more beneficial conditions for the corundum grain growth may
occur and that may result in a decrease of hardness. As expected, the ceramics
manufactured with a higher content of the starting alumina powder with a lower
particle size and a higher specific surface have more uniform microstructure and,
as a sequence, higher hardness and other physical properties. Also, the ceramics
with a smaller grain size have a narrower standard deviation in hardness. The
maximal hardness values (HV10 greater than 1550 kg/mm2) are observed for the
AL99.6 ceramics and for the AZ ceramics, which have a very uniform
microcrystalline microstructure with practically no glassy phase.
Indentation fracture toughness of the studied alumina ceramics tends to
increase with the Al2O3 content like Young’s Modulus and hardness; however,
this rise is not significant. As expected, the presence of the zirconia phase in the
alumina matrix results in an increase of fracture toughness.
Slip cast alumina ceramics demonstrate higher values of mechanical
properties such as flexural strength, hardness, Young’s Modulus and sonic
velocity, than pressed ceramics. It is explained by a higher level of densification
and uniformity and less stress and fewer defects formed during slip casting and
binder-burn out stages.
Ballistic Performance
The fracturing process of ceramics during impact and penetration at bullet
velocities ranging from 700 to 5000 m/s has several stages, and it includes [3]:
1) initial impact with hydrodynamic flow of penetrator and armor ceramics;
2) breakup and continued flow of penetrator and high speed jetting of debris;
3) ceramic fracture, formation of Hertzian cone cracks, and tensile cracks on the
back face with continued penetrator breakup and flow;
4) erosion of penetrator and widespread fracture of ceramics.
With increasing bullet velocities, the energy transmission through ceramic
armor and across boundaries via shock waves becomes more valuable, i.e. the
ability of ceramics to dissipate the bullet kinetic energy and to prevent the crack
propagation is very important. Energy dissipation during bullet impact and
96 Ceramic Armor Materials by Design
fracturing of ceramics depends on many factors dealt with ballistic situation
(including initial bullet kinetic energy, bullet material properties, etc.) and
properties and microstructure of ceramics. Regarding ceramics, it should have
some level of properties, which include density and porosity, hardness, fracture
toughness, Young’s Modulus, sonic velocity, mechanical strength, and some
others. Any single property does not have a direct correlation with ballistic
performance because the fracture mechanism during the bullet impact is very
complicated, the crack formation is caused by different stress factors and it occurs
in an extremely short time. In short, the microstructural features affecting physical
and ballistic properties strongly influence crack propagation and energy
dissipation mechanisms and ultimately ballistic performance. Hence, all relevant
properties, as well as ceramic microstructural features, must be carefully
considered in assessment of ballistic performance of protective systems.
For dense homogeneous armor alumina ceramics in order to achieve
acceptable consistent level of ballistic performance Vickers hardness HV10
should exceed 1220-1250 kg/mm2 (i.e. significantly exceed the projectile
hardness). Sonic velocity indicating the ability of hard ceramics to dissipate
energy from the impact area should be greater than 10,000 m/s (preferably,
10,500-11,500 m/s). Young’s Modulus should be greater than 325 GPa (usually
350-450 GPa dependent on the Al2O3 content). The impedance I= c=( E)1/2
[3]
(where is density, c is sonic velocity, E is Young’s Modulus) indicating a wave
propagation in a material should have a level similar to steel (400 MPa.s/m).
Flexural strength should be greater than 220 MPa. Although many authors [1-3,7]
indicate that armor ceramics should have low fracture toughness, it seems that KIc
should not be lower than 3 MPa.m0.5
. Some “balance” between levels of hardness
and fracture toughness needs to be maintained.
There were a number of attempts to describe ballistic performance using
mathematical modeling (e.g. [1-4,7]). All of these had different approaches but
they did not fully describe ballistic performance. However, these models help to
understand better the fracturing mechanism, indicate the important properties
relevant to ballistic performance and allow for a preliminary evaluation of
ballistic performance. For example, Neshpor, et al., 1995 [7] proposed the semi-
phenomenological criterion of evaluation of ability of ceramics to dissipate
ballistic energy using the formula: D = 0.36 (HVcE)/KIc2. The approximate values
of the ballistic energy dissipation criterion have been calculated (Table 1). This
formula and the calculated D-criterion values show that the highest hardness, or
the lowest fracture toughness, is not the dominant factors affecting ballistic
performance. The optimal combination of relevant factors should be considered
for the promising armor ceramics. As an example, the AZ ceramics with high
hardness and elevated fracture toughness (in comparison with alumina ceramics)
demonstrate high ballistic performance.
Ceramic Armor Materials by Design 97
All the studied ceramics demonstrate a high level of ballistic performance.
The armor systems based on these alumina ceramics bonded with appropriate
aramid-based and fiberglass backing materials are capable of defeating 7.62x51-
mm and 7.62x63-mm AP FMJ ammunition and 7.62x39-mm and 7.62x51-mm
Ball FMJ ammunition, and they provide ballistic protection to Level III or Level
IV dependant on the ceramics and backing material thickness (Level IV in
conjunction with a ballistic vest). The armor systems for personnel protection
have satisfactory multi-hit ballistic performance (up to 6 hits to one body-armor
plate). Trauma for the armor plates for personnel protection made from these
materials occurred at acceptable levels (i.e. not greater than 44-mm deformation
in accordance with NIJ Standards). The alumina ceramics with higher hardness
demonstrated less trauma and bullet intrusion. However, in this case, a greater
degree of a crack growth is observed, probably, due to a higher “ratio” between
hardness and fracture toughness. As mentioned above, a bullet is distorted and
eroded during the initial contact with the ceramics; the erosion of the projectile is
greater as hardness of ceramics increases. As expected, the highest level of bullet
erosion was observed for the hardest ceramics such as alumina-zirconia AZ and
alumina AL99.6 and AL98.5 ceramics that correlates well with the data [6].
Different kinds of cracks are formed during the ballistic impact, which depend
on type of a bullet and properties of ceramics. A locus of conoid coaxial cracks
starts at the impact point; radial tensile cracks are initiated at the back surface
close to the axis of impact. Star cracks are formed at the side of conoids.
Tangential spall cracks occurred due to shear stress waves reflected from the
edges of a tile and due to the formation of the cone cracks; lateral spall cracks
may also form due to longitudinal stress waves reflected from the backing
support. Comminution and erosion of ceramics occur at the cone area. The
thickness of ceramics may also affect crack formation and development. Usually
more conoids occur with greater thickness. The nature and thickness of backing
materials (high-strength aramid-based fabric such as Kevlar, aluminum sheet,
polymer block, or others) may have a significant influence on crack propagation
due to their different abilities to reduce the stress. Fragments of damaged
ceramics with different sizes ranging from big chunks to a fine powder were
observed after fracturing. The chunks with bigger sizes were observed for the AZ
ceramics and for the ceramics with a relatively lower content of crystalline
phases, such as AL97ML and AL98. By contrast, dense boron carbide and silicon
carbide armor ceramics commonly demonstrate explosive shattering at the
shooting. This transforms to a powder at the damaged area with a minimum
amount of large ceramic chunks, which do not remain with the backing material
despite a high level of bonding.
The compaction of the comminuted ceramics under the compression resulting
from impact with a projectile affects penetration resistance. The comminuted and
98 Ceramic Armor Materials by Design
compacted area of the studied alumina ceramics had relatively small chunks and
agglomerated powder. Microscopic observation showed that the ceramics had
multi-grain fragments ranging from 10 to 100 µm, with micro-cracks and elevated
porosity. In contrast, the uncomminuted area had only macroscopic cracks. The
alumina ceramics with a higher content of a glassy phase such as AL97ML and
AL98 had more chunks; ceramic fragments had more micro-cracks, which
develop mostly through a glassy phase. The fine-crystalline ceramics with an
insignificant amount of a glassy phase had fewer chunks; micro-cracks grew
through the grain boundary and even through grains. Some grains were pulled out
which resulted in elevated porosity. The AZ ceramics had more chunks, but
micro-cracks with a relatively short length grew through the grain boundaries.
Considering armor systems with satisfactory ballistic performance, ceramic
damage at impact should be more conical than cylindrical. At the same time, the
hole caused by a bullet should have a small size. This indicates that the bullet
velocity decreases significantly after the contact with hard ceramics and, hence,
trauma should be minimized. Cracks in the ballistically tested ceramics are
desired to be shorter with small cones. In this case, the residual part of a
ballistically tested ceramic plate will have less damage, and, therefore, a ceramic
plate used for personnel protection has a better probability to resist subsequent
shots. This is true when a ceramic and a backing material are still bonded after
shooting; i.e. the adhesive and the bonding technique are optimized. Alumina
ceramics with different compositions and structure may be used for particular
ballistic applications. Ceramics with a higher alumina and less glassy phase
content and higher hardness are more beneficial for armor tile manufacturing and
for single-hit applications. Ceramics with a higher glassy phase content and lower
hardness values (or, probably more correct, with a lower hardness/fracture
toughness “ratio”), are more suitable for multi-hit ballistic applications despite
possibly demonstrating greater trauma. However, these recommendations are
broad generalizations, and again, all relevant properties, including the ability to
dissipate the bullet energy, must be considered for complete analysis of required
ballistic properties. The curvature of monolithic armor plates may affect the
fracturing, macro-crack propagation, and multi-hit performance.
SUMMARY
The developed and studied alumina ceramics with an Al2O3 content of 97-
99.6-wt.%, as well as the alumina-zirconia ceramics, demonstrate a high level of
physical properties and high ballistic performance. High performance of these
ceramics is achieved by maintaining the proper composition, including the use of
raw materials with optimal parameters, and microstructure, as well as through
optimization of the manufacturing process and quality control points. Properties
affecting ballistic protection and ballistic test results are discussed. A combination
Ceramic Armor Materials by Design 99
of relevant properties for ballistic protection, including microstructure features
should be considered in the evaluation and selection of ceramics used in armor
applications.
REFERENCES
[1]. C.F. Cline and M.L. Wilkins, “The Importance of Material Properties in
Ceramic Armor”; p.p. 13-18 in DCIC Report 69-1; Part I: Ceramic Armor, 1969.
[2]. Soon-Kil Chung, “Fracture Characterization of Armor Ceramics”, Amer.
Ceram. Soc. Bul., 69 [3] 358-366 (1990).
[3]. D.J. Viechnicki, M.J. Slavin, and M.I. Kliman, “Development and Current
Status of Armor Ceramics”, Amer. Ceram. Soc. Bul., 70 [6] 1035-1039 (1991).
[4]. I.Yu. Kelina and Yu.I. Dobrinskii, “Efficiency of the Use of Silicon
Nitride Ceramics as an Armor Material”, Refractories and Technical Ceramics (in
Russian), [6] 9-12 (1997).
[5]. B. Matchen, “Application of Ceramics in Armor Products”; p.p. 333-342
in Advanced Ceramic Materials; Ed. by Hamid Mostaghasi; Key Engineering
Materials, Vol. 122-124, 1996. Trans. Tech. Publications, Switzerland.
[6]. R.G. O’Donnell, “An Investigation of the Fragmentation Behaviour of
Impacted Ceramics”, J. of Materials Science Letters, [10] 685-688 (1991).
[7]. V.C. Neshpor, G.P. Zaitsev, E.J. Dovgal, et al., “Armour Ceramics
Ballistic Efficiency Evaluation”; p.p. 2395-2401 in Ceramics: Charting the
Future. Proc. 8th
CIMTEC Florence, 28 June-4 July 1994; Ed. by P. Vincenzini,
Techna Srl., 1995.
100 Ceramic Armor Materials by Design
Table 1. Some physical properties of the studied alumina and alumina-zirconia armor ceramics
Property AL97ML AL98 AL98.5 AL99.6 AZ
Density,* g/cm3
3.74- 3.76 3.78-3.82 3.81-3.84 3.90-3.91 4.35-4.39
Young’s Modulus, GPa 280-300 325-360 370-420 400-450 310-340
Sonic Velocity, km/s 9.5-9.9 10.0-10.5 10.6-11.3 10.7-11.6 9.8-10.0
Vickers Hardness HV10, kg/mm2
1230-1260 1250-1330 1320-1420 1520-1560 1520-1580
Fracture Toughness KIc, MPa.m1/2
3.0-3.3 3.2-3.3 3.3-3.4 3.1-3.4 3.9-4.0
Flexural Strength, MPa - 250-350 270-360 320-380 500-560
Ballistic Energy Dissipation Criterion
Dx10-12
, 1/s (calculated)
1.70-1.95 1.50-1.60 1.80-1.95 2.20-2.40 1.15-1.20
* Water absorption is not greater than 0.02%
These data are performed for the materials manufactured by slip casting and pressing
Ceramic Armor Materials by Design 101
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BALLISTIC PERFORMANCE OF ALUMINA CERAMIC ARMORS
Murat Vural and Zeki Erim B. A. Konduk
Istanbul Technical University Bogazici University
Dept. of Aeronautics Institute of Biomedical Eng.
Maslak-Istanbul 80626 Dept. of Materials
Turkey Bebek-Istanbul 80815
Turkey
A.H. Ucisik
Bogazici University
Institute of Biomedical Eng.
Dept. of Materials
Bebek-Istanbul 80815
Turkey
ABSTRACT
High quality alumina ceramic tiles, backed with semi-infinite aluminum
blocks were ballistically tested with armor piercing 7.62 mm projectiles. The
failure mechanism, ballistic efficiences and fragmentation behavior of ceramics
were investigated under impact loading conditions. The thickness and projectile
velocity were essential. Ballistic efficiency was affected by the thickness and
projectile velocity. Upon impact, radial, cone and lateral cracks form and
disintegrate the ceramic tile. A ceramic conoid zone within the innermost cone
crack interacts with the projectile. In the present study, ballistic efficiencies that
quantify the normalized performance of ceramic against the impacting projectile,
have been found to be a function of the projectile velocity and ceramic tile
thickness. High values of ballistic efficiencies were achieved for thinner ceramic
tiles and for higher impact velocities. These effects of projectile velocity and
ceramic tile thickness on the ballistic efficiency are thought to be extremely
important when making merit ratings between armor ceramics tested at various
velocities or thickness.
Ceramic Armor Materials by Design 103
To the extent authorized under the laws of the United States of America, all copyright interests in this publication are the propertyof The American Ceramic Society. Any duplication, reproduction, or republication of this publication or any part thereof, withoutthe express written consent of The American Ceramic Society or fee paid to the Copyright Clearance Center, is prohibited.
INTRODUCTION
Impact, penetration, perforation and trauma effects imposed on materials,
including living materials, are important to many fields of engineering, including
biomedical engineering, orthopedic and traumatology and to the military for
armor application. This interest comes from the desire of both increasing the
penetration capability of projectiles and making protective armor systems
resistant to certain types of threats. Penetration may be defined as the entrance of
a projectile into a target without completing its passage through the body.
Perforation, on the other hand, implies the complete piercing of a target by the
projectile. Upon striking of a projectile, a target can fail by a variety of
mechanisms, depending on a long series of parameters such as impact velocity,
geometry of interacting bodies, and material properties of both the projectile and
the target.
Armors are a means of protection against penetrators. The item to be protected
may be a human body, a vehicle or a fixed building. Body armor, the most
important one, is intended to protect individuals primarily against fragments from
high-explosive artillery shells, grenades, fragmenting mines, as well as projectiles
from small arms and rifles.
The evaluation of ballistic performance is one of the most important issues in
the selection of an armor material. However, the measurement of ballistic
performance has always been a very difficult task because of the destructive
nature of ballistic testing and the many variables involved, such as the type and
velocity of threat, and the types of target and target support.
A method to determine the ballistic performance that found early acceptance,
especially for small arms threats, is the V50 test. In this test, the efficiency of the
tested ceramic tile is determined by the magnitude of the ballistic limit velocity
(VBL), defined as the impact velocity at which 50 percent of the projectiles do not
penetrate the target. The experimental target configuration for V50 testing consists
of bonding a ceramic tile to a backup plate of comparable thickness and shooting
projectiles at these targets. However, as Rosenberg and Yeshurun [l] pointed out,
the V50 test with a thin-backing configuration is not a good test for screening the
ballistic performance of brittle materials.
Another method commonly used to evaluate the ceramics for armor
applications is to fire a reference shot into a thick reference backup target and a
second shot through the candidate ceramic tile which is bonded to the same
backing material, and to compare the residual depths of penetration. The
application of this ballistic testing technique [1-4], i.e. so-called thick backing
technique, has seen the use of various projectile types, backing types and
conditions of lateral confinement. The thick backing technique, originally
introduced by Rosenberg et al. [see Ref. l] and shown schematically in Figure la,
104 Ceramic Armor Materials by Design
also enables a convenient measure for the ballistic efficiency ( ) of ceramic tiles,
which can be expressed as
= CC
RBB
t
P-P (1)
where B and C are the densities of the backing (aluminum in the present study)
and ceramic respectively, tC is the ceramic thickness, and (PB -PR) is the reduction
in thickness of backing penetrated due to the ceramic tile being in place, i.e. the
difference between the reference depth and the residual depth.
In the present work, ceramic tile thickness and projectile velocity are altered
in order to determine their influence on the ballistic efficiency parameter. A prime
objective is to understand the reliability of the ballistic merit ratings based on
thick backing technique.
EXPERIMENTAL STUDY
The ceramics used included two grades of alumina (prefix AD) and are listed
with their physical and mechanical properties in Table 1. The ceramic tiles were
50 mm square and of six different thicknesses, ranging between 4.1 and 14.7 mm.
These tiles were provided from Kaleporselen A.S., Istanbul, and their listed
properties were taken from the previous study of Birbilen et al. [5].
Table I. Properties of ceramic plates [5].SinteringCeramic
Plate Temp (°C) Time (h)
Density
(gr/cm3)
Hardness
(GPa)3-Point Bending
Streng. (MPa)Compressive
Streng.(MPa)
AD-96 1650 3 3.80 14.5 360 1460
AD-99.8 1680 3 3.90 15.0 400 1600
The ceramic tiles were bonded to clean 6061-TO aluminum alloy backing
blocks using a neoprene-based adhesive. Two configurations of target setup,
referred to as ''thick backing'' method, are schematically shown in Figure 1a and
1b. In all cases, the targets were impacted by a 0.30 calibre (7.62 mm) conical-
nosed armor piercing projectile at velocities ranging from 576 to 803 m/s. A
schematic of the projectile cross section is shown in Figure 2. The projectile has a
total mass of 9.56 0.08 g. This mass contains a 3.5 g hardened steel penetrator
with a conical nose.
Ceramic Armor Materials by Design 105
Figure 1. Schematics of measured penetration depths; (a) in bare aluminum
blocks, (b) in target panels with "ceramic tiles (thick backing
configuration). The depths were measured by x-ray radiography and/or
depth gauging measurements.
Figure 2. Schematic of the projectile cross section
showing the steel core and jacket.
EXPERIMENTAL RESULTS AND DISCUSSION
Experimental data is presented in Table 2, in the form of reference depth into
the backing plate, residual depth into the backing following perforation of the
ceramic tile, and ballistic efficiency parameter, , as defined by Eqn.1. The
results presented in Table 2 are for successful shots on ceramic tiles smaller than
8.3 mm in thickness; ceramics of greater thickness could not be perforated by the
projectiles in the velocity range investigated. Therefore, ballistic efficiency for
those thicker tiles is not calculated since it is not appropriate. In some cases, a
greater sign (>) is shown just prior to value. These tiles were too thick for the
7.62 mm AP projectile and there was no penetration into the aluminum backing.
Thus, a lower bound on the ballistic limit can be calculated as if the ceramic tile
were just perforated.
Table 2. Penetration data of AD-99.8 and AD-96 ceramic tiles.
106 Ceramic Armor Materials by Design
Specimen
Code
Projectile
Velocity
(m/s)
Residual
Penetration
(mm)
Reference
Penetration
(mm)
Ballistic
Efficiency
( )
X9914.15/4 795 46 95.3 8.4
X9914.1/2 803 42 96.6 9.5
X9914.1/3 797 33 95.6 10.9
X9914.1/5 788 34 94.2 10.4
X9914.1/6 672 20 75.5 9.6
X9914.1/8 601 21 64.1 7.5
X9916.5/1 802 8 96.4 9.7
X9916.4/2 802 6 96.4 10.0
X9916.4/3 785 5 93.7 9.8
X9916.5/5 694 3 79.1 8.3
X9916.5/6 680 3 76.8 8.1
X9916.45/8 585 1 61.5 6.7
Y9918.0/1 800 2 96.1 8.4
Y9918.0/3 800 4 96.1 8.2
Y9918.2/4 793 4 95.0 7.9
Y9918.1/5 692 1 78.7 6.8
Y9918.0/2 803 5 96.6 8.1
Y9918.1/7 615 0 66.4 >5.8
Y9918.3/8 595 0 63.1 >5.4
Y9918.1/6 693 1 78.9 6.8
X9614.2/7 595 16 63.1 8.0
X9614.2/2 792 33 94.8 10.5
X9614.2/3 803 37 96.6 10.1
X9614.2/6 659 23 73.4 8.5
X9616.0/7 592 8 62.7 6.5
X9616.0/3 804 11 96.7 10.2
X9616.0/1 804 9 96.7 10.4
X9616.0/2 795 6.5 95.3 10.5
X9616.0/4 789 7.5 94.3 10.3
X9616.0/8 577 4 60.2 6.7
X9616.0/6 677 6 76.3 8.3
Y9618.3/3 785 1 93.7 7.9
Y9618.3/4 778 5 92.6 7.5
Y9618.3/6 658 4 73.3 5.9
Y9618.3/2 800 0 96.1 >8.2
Y9618.3/5 678 0 76.5 >6.5
Y9618.3/7 592 0 62.7 >5.4
Y9618.3/8 576 0 60.1 >5.1
Ceramic Armor Materials by Design 107
Table 2. Penetration data of AD-99.8 and AD-96 ceramic tiles.
Figure 3 shows the variation of ballistic efficiency data in Table 2 against the
projectile velocity for the ceramics of various thicknesses. These data suggest that
ballistic efficiency of a ceramic tile depends on both the impact velocity and the
thickness of ceramic tile used. It is well established by previous workers that the
ballistic efficiency of ceramics is affected by the kind of ceramic used [1] and the
projectile geometry [4]. The results of the present study (see Figure 3) reveal that
the thickness of the ceramic and the velocity of projectile are two new factors that
significantly effect the ballistic efficiency ( ) of ceramics. Increased projectile
velocity or decreased tile thickness results in increased ballistic efficiency. This
result is of considerable importance because it suggests that one must be very
careful when making merit ratings based on thick backing technique for ceramics
of different thickness (which is the case in Ref. [1]) and for ceramics impacted at
different velocities (which is the case in Ref. [4]).
Rosenberg and Yeshurun [1] compare the ballistic efficiencies of different
types of ceramics such as SiC, B4C, TiB2 and Al2O3 by using the thick backing
technique. However, the thickness of the ceramics in their work ranges between 6
and 10 mm. In the light of the results of present study, this type of a comparison is
open to speculation. Findings of the present study suggest that a ceramic tile of 6
mm will show greater ballistic efficiency than the 10 mm tile. Therefore, a
comparison between the types of ceramics without regarding the effect of
thickness will result in unfair merit ratings, which are based on thick backing
technique and Eqn.1.
In a similar manner, Woodward and Baxter [4] investigated the effect of
projectile geometry on the ballistic efficiency of Al2O3 ceramics. They also used
the thick backing technique in their experimental setup. Even though they held the
ceramic tile thickness constant, the projectile velocity ranged between 899 and
1243 m/s in their study. The results of present study show that increased projectile
velocity produces increased ballistic efficiency for Al2O3 ceramics in the range
between 576 and 803 m/s. If this trend holds for the range between 899 and 1243
m/s, some of the comparisons made in the work of Woodward and Baxter [4]
seem to be open to speculation as in the work of Rosenberg and Yeshurun [1].
CONCLUSIONS
Terminal ballistic tests were performed on high quality alumina ceramic tiles
backed with thick aluminum plates, i.e., the so-called thick-backing method. The
two main parameters were the projectile velocity and the thickness of ceramic
tiles. Results clearly show that the ballistic efficiency parameter of ceramics is not
constant and, contrary to common assumption in scientific literature, it
significantly varies as a function of both the tile thickness and the projectile
velocity, at least for the range of velocities between 576 and 803 m/s and
thickness between 4.1 and 8.3 mm. These effects of velocity and thickness on the
108 Ceramic Armor Materials by Design
ballistic efficiency are thought to be extremely important when making merit
ratings among the armor ceramics tested at various velocities or thickness.
Figure 3. Ballistic efficiencies (defined by Eqn.3.1) of ceramic tiles for varying
projectile velocities: (a) AD-99.8 ceramics, (b) AD-96 ceramics.
REFERENCES1
Rosenberg, Z. and Yeshurun, Y., “The Relation Between Ballistic
Efficiency and Compressive Strength of Ceramic Tiles”, Int. J. Impact Engng.,
Vol. 7, No.3, 357-362, (1988).
Ceramic Armor Materials by Design 109
2 Rosenberg, Z., Bless, S.J. and Brar, N.S., “On the Influence of the Loss of
Shear Strength on the Ballistic Performance of Brittle Solids”, Int. J. Impact
Engng., Vol.9, No.l, 45-49, (1990). 3 Rosenberg, Z. and Tsaliah, J., “Applying Tate's Model for the Interaction
of Long Rod Projectiles with Ceramic Targets”, Int. J. Impact Engng., Vol.9,
No.2, 247-251, (1990). 4 WOODWARD, R.L. and BAXTER. B.J., “Ballistic Evaluation of
Ceramics: Influence of Test Conditions”, Int. J. Impact Engng., Vol.15, No.2,
119-124, (1994). 5 Birbilen, M., Yildirim, I. and Valenta, L., Introduction of High-Tech
Ceramics into Turkish Industry (in Turkish), in M.L. Ovecoglu & H. Yaparlar
(eds), Proc. 2nd
Int. Ceramic Congress, Istanbul, 24-28 Oct. 1994, Turkish
Ceramic Society Press, Istanbul, Vol.2, 469-474, (1994).
110 Ceramic Armor Materials by Design
Penetration and Ballistic Testing
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AN OVERVIEW OF BALLISTIC TESTING METHODS OF CERAMIC
MATERIALS
Dr. Michael J. Normandia
Armor Mechanics Branch,
Terminal Effects Division,
Weapons & Materials Directorate,
Army Research Laboratory
AMSRL-WM-TA
Aberdeen Proving Ground
APG, MD 21005
Mr. William A. Gooch
Armor Mechanics Branch,
Terminal Effects Division,
Weapons & Materials Directorate,
Army Research Laboratory
AMSRL-WM-TA
Aberdeen Proving Ground
APG, MD 21005
ABSTRACT
An overview of impact testing techniques used to characterize or evaluate
engineering structural ceramics for armor applications is presented. The required
brevity of this paper restricts the detail to a listing of the commonly used testing
methods, a categorization of ballistic techniques, and an extensive, but far from
complete, listing of key references appears in alphabetical order, and we
apologize for any omissions. Individual speakers have been invited to this
conference, and they will provide greater detail of the testing techniques, the
evaluation procedures, and of the armor system design methodologies. In
addition, the presentation provides typical testing configurations, typical results,
and test objectives. A similar presentation and companion paper on ceramic
armors by Gooch also provides information on how this data is typically used to
construct armor systems.
A categorization of the testing techniques is provided to classify testing
methods into those that attempt to characterize a particular ceramic material’s
properties, those that attempt to evaluate and rank a ceramic material’s
performance for armor applications, and those techniques that actually evaluate
the ceramic in configurations more representative of armor systems. Finally, we
discuss some of the difficulties in utilizing these testing techniques for ranking
ceramic materials, particularly due to the fairly recent discovery of dwell, which
has had profound effects on data evaluation. Dwell describes the behavior of an
eroding penetrator prior to penetration and in certain circumstances a penetrator
can completely erode on the ceramic surface without penetration.
Ceramic Armor Materials by Design 113
To the extent authorized under the laws of the United States of America, all copyright interests in this publication are the propertyof The American Ceramic Society. Any duplication, reproduction, or republication of this publication or any part thereof, withoutthe express written consent of The American Ceramic Society or fee paid to the Copyright Clearance Center, is prohibited.
EDITORIAL NOTEA subset of the presented vuegraphs appears in a separate appendix in this manuscript.
These are presented in a Powerpoint report format beginning on page 131.
INTRODUCTION
During the past quarter-century, ceramics have seen limited use in armor
applications, mostly for lightweight armor applications such as body armor,
helicopter seats, and appliqués on land vehicles for additional threat protection.
Renewed interest and an increase in applications have occurred recently due to the
emphasis on weight reduction. Numerous ballistic testing techniques have been
used to ascertain the effectiveness of ceramics in armor applications, and, in
particular, to rank the performance of the various candidate materials. Typical
impact threats are fragments, which are representative of exploding warheads,
soft-core bullets, hard-core bullets, medium to long-rod, kinetic-energy
penetrators, and high-velocity, shaped-charge jets. Different defeat mechanisms
dominate each of these different threats, and for each threat the dominant ceramic
failure mechanism in a typical armor system is also likely to be different.
The ranking of ceramic (and other brittle) materials is significantly
affected by the test configuration’s geometry. Since armor designers utilize
different armor configurations to exploit different threat defeat mechanisms, the
ranking of ceramic materials is also clearly threat dependent. These are
significant complicating factors, which affect the results of even fundamental
testing methods, as the particular choice of threat and geometric configuration
may affect the ranking. Thus, evaluation of a material for a particular application
may require a test that best represents the mechanisms associated with that
application.
The influence of target, or test configuration geometry, is primarily due to
the material’s very strong pressure dependent strength (in both intact and
damaged states), and the very weak tensile failure strength. Geometry affects the
response to the initial shock, confining pressure, and greatly affects the onset of
tensile failure. Tensile failure may occur for a variety of different reasons, such
as local shear deformation, global ceramic bending, reflected tensile waves from
boundaries or free surfaces, etc. Once failed, the ceramic strength, depends very
strongly on the confinement (pressure) and the fragments can easily be displaced
from the penetrator path, sometimes providing very little resistance to penetration,
if the geometry permits. These effects introduce different time scales into the
testing, such as the time to fracture, the propagation of damage, and the time
response of the containment system, relative to the penetration time. The
outcome of a particular test depends upon when a particular failure mechanism
occurs, in both an absolute and a relative sense, which depends upon the
particular materials used, making comparisons relative rather than absolute. A
change of impact conditions, or material thicknesses may alter the ranking.
These are the most likely reasons why the numerous attempts to correlate
actual armor ballistic performance to fundamental material properties have been
114 Ceramic Armor Materials by Design
unsuccessful. However, partial success in ranking materials ‘potential’ has been
made using a variety of techniques where a particular class of threats and a
common defeat mechanism were used in the testing technique, and where the
geometry was carefully controlled.
TEST CATEGORIZATION
For simplification (and not for completeness), the types of testing can be
categorized, subjectively, as either phenomenological, armor-material
characterization, or armor-design oriented.
Phenomenological experimental methods attempt to determine or obtain
specific material properties and include shock physics impact data, such as wave
profiles in normal or oblique plate impact. These properties are then utilized to
evaluate a material’s potential performance in an armor system, often utilizing
numerical simulations. They are non-ballistic and in non-armor configurations,
but are important and essential tests, particularly for constructing constitutive
models used in numerical simulations for armor design.
Armor-material characterization experimental methods, attempt to
determine a particular ceramic material’s resistance to penetration (often in an
integrated sense), and include the traditional depth-of-penetration experimental
technique. These tests are dynamic and are sub-divided into non-traditional tests,
with border on phenomenological, and traditional ballistic testing methods, which
are often used to validate numerical simulations. It is unfortunate, that at this time,
these test results are necessary to develop constitutive models, mainly to
determine the strength of the damaged ceramic.
Armor-design oriented experimental methods include the traditional, MIL-
SPEC ballistic limit velocity experiments, to determine a v50, for example. This
category also includes armor design testing methodologies to help isolate an
optimum armor configuration to defeat a specific threat.
A listing of the techniques in each of the three categories appears below,
followed by a brief description of the objectives. The experimental methods can
be (subjectively) categorized as either phenomenological (Table I), armor-
material characterization (Table II), or armor-design oriented (Table III). Ballistic
experimental techniques are bolded and will be discussed in slightly more detail.
Phenomenological Experiments
All of the experimental techniques listed in this category are not ballistic
testing methods, but they are necessary to determine the fundamental material
properties or behavior under shock loading. These properties and behavior are
utilized to develop constitutive models or to validate numerical simulation tools.
The often-attempted goal of ballistic ceramic performance testing methods is to
relate the performance back to these more fundamental characteristics. This has
Ceramic Armor Materials by Design 115
been partially successful, but not generally recognized or utilized. These tests can
also be considered to be material characterization tests in that they characterize a
particular material property.
Table IA Categorization of Testing Techniques – Phenomenological Experiments
CATEGORIES OF CERAMIC TESTING UTILIZED TO EVALUATE
ARMOR PERFORMANCE OF CERAMICS
Phenomenological Experiments
(1) Pressure-volume
(2) Plate impact (normal, oblique, or multiple impacts-reshock)
(3) Split Pressure (Hopkinson or Kolsky) Bar (Compression)
(4) Bar impact (typically bar impacting bar)
(5) Tensile or Torsion
(6) Quasi-static three- or four-point bending tests
(7) Quasi-static indentation
A combination of these fundamental test results, provides significant
information, which, when utilized with numerical simulations, provides a
realizable-hope that ranking can be obtained with minimal testing. These
methods typically provide information about the best achievable performance.
One of the most promising is the indentation techniques (static or dynamic) to
measure stress-strain relations, dynamic yield strength, and apparent plasticity.
There are a variety of other phenomenological testing methods used that
do not readily fit into any of these categories. However, while this categorization
is likely to be incomplete, most traditional tests are included. Perhaps more
importantly, the variety of categories and testing methods demonstrates the
variety of experimental techniques and data that are typically generated for a
single processed ceramic material. Further complications are due to variations
attributed to the starting powder source and impurities, the batch processing
technique and the particular manufacturer. Additional differences may be due to
non-uniformities within a large processed sample cut up into smaller samples,
surface preparation, ceramic geometric configuration, and geometric testing scale.
Improved testing techniques have been developed that take advantage of
numerical simulation tools to design a test to prevent failure before the
measurement and to identify a particular, desired stress-state. Two examples
were demonstrated by Sandia National Labs to examine impact of confined
cylindrical ceramics at high strain rates using graded impactors (Chhabildas) and
to provide a ramp-loading time history in Split-Hopkinson Pressure Bars at
intermediate strain rates (Forrestal, Frew).
116 Ceramic Armor Materials by Design
Underlying the numerical simulations are constitutive models, which
describe the material behavior in both a fully intact state, and in a damaged state,
both of which have pressure dependent material strength. The behavior of
ceramics in an intact, damaged or failed state, particularly under high strain-rate
and high-pressure loading, typical of ballistic impact events, has led to
constitutive models that are empirically determined. Typical formulations require
the use of ballistic data to calibrate the model coefficients, in particular, the
criterion for the transition from an intact to a damaged or failed state, and the
strength of the partially or fully-damaged material. In addition, many models
accumulate damage and the partial damage states often degrade particular
material properties (e.g., moduli), for which no data exists. Additional testing
techniques attempt to characterize the failed material in non-ballistic experiments,
such as collapsing cylinders or spherically expanding cavities. The connections
between cavity expansion and penetration have made this class of experiments
very relevant to material performance as well.
Armor-Material Characterization Experiments
This category of testing methods utilizes dynamic impact. These tests are
also phenomenological in nature, but these tests are utilized to directly measure,
or determine from behavioral models, properties characteristic of target resistance
or resistance to penetration. These experiments typically control the geometry of
the test, although variations exist between testing agencies. These tests were
specifically developed to attempt to directly or indirectly evaluate, rank and or
compare ceramic performance for ballistic armor applications, due to the inability
to utilize the phenomenological experimental data for this purpose. The recent
addition of the dwell/penetration transition experiments were added to this
category, even though below some impact load (velocity) there is no penetration.
Table II Categorization of Testing Techniques – Armor-Material
Characterization Experiments
CATEGORIES OF CERAMIC TESTING UTILIZED TO
EVALUATE ARMOR PERFORMANCE OF CERAMICS
Armor-Material Characterization Experiments Ref(CEX) Cavity expansion or cylindrical collapse 11, 52-4
(DAM) Damage Propagation (edge on impact) 31
(IND) Indentation: dynamic or loading and unloading 37, 38, 56
(NDP) Non-deforming penetration (referred to as rigid penetration) 10, 12, 49
(PEN)Semi-infinite penetration vs. velocity time histories 21, 42-4, 57
(DOP)Modified depth-of-penetration experiments (quantifying dwell) See Table IV
(DWE) Complete dwell (for damage onset and for structural response) See Table IV
(DPT) Dwell/penetration transition (concerns about shock mitigation) See Table IV
Ceramic Armor Materials by Design 117
Non-Traditional Experiments-The following four test techniques are non-standard
techniques used as an extension of the phenomenological models to provide
material properties directly applicable to the prediction of the performance in
armor systems.
CEX: Cavity Expansion or Cylinder Collapse: Cavity expansion models
have been successfully utilized in penetration modeling for brittle materials,
typically geologic materials, such as soils and limestone, and concrete. Extensive
utilization of these models for ceramics has also achieved significant progress.
The basic premise is that the pressure at the penetrator/target interface, which
provides the resistance to penetration, can be computed from integration of the
equations of motion over the entire target. Hence the quantification of the various
damage regions, brittle comminuted, brittle fractured, plastic, and elastic are
necessary to obtain this information. In addition, the theories have been extended
to account for the dynamic evolution of the cavity and the various regions and
account for the presence of finite boundaries as well.
DAM: Damage Evolution Experiments: The consequences of damage can
be described in several aspects. First, for thin tiles that may fail in tension due to
bending, for example, less resistance to penetration is provided, even though the
material is essentially still capable of providing a compressive stress (particularly
under confining pressure). In thicker tiles, or those subject to more lethal threats,
the extensive fracturing and comminution occurs early in the process, hence most
of the penetration occurs in damaged material, which is also a function of
confining pressure. Lastly, the evolution of a damage front has been identified,
and if a threat were to penetrate faster than this damage front, greater resistance to
penetration will be provided. Extensive data on shaped-charge impacts have
quantified an order of magnitude increase, and is presented in the references.
Edge on impacts have attempted to quantify the propagation velocity of the
damage front. Discussions in the literature refer to a damage wave, but that
concept is debatable and the subject of current research. The reference describes a
test technique to monitor the damage front from an edge-on impact.
IND: Indentation: Quasi-static and dynamic indentation experiments
typically are used to measure hardnesses of materials and are representative of the
materials compressive strength, a logical first property to examine when
attempting to rank materials. However, brittle materials resistance to indentation
change when they are cracked or fractured, and typical indenters create localized
shear that can make the measurements difficult to interpret. Hertzian contact has
also been used extensively (Lawn). The additional information provided by the
use of multiple indents is the generation of stress-strain curves, which yield
information about the transition from an elastic behavior to an inelastic one. The
behavior after this transition is equally important in the ceramics ability to resist
penetration. The comparison of the apparent plasticity to the dynamic yield
118 Ceramic Armor Materials by Design
strength has been used to rank ceramics. These techniques are promising in
providing valuable, fundamental material properties, that potentially can be used
to quantify ceramics (particularly when combined with so-called first-principle
numerical simulation tools.
NDP: Non-Deforming Penetration: Non-deforming or rigid penetration
has been used by numerous researchers to measure the resistance to penetration of
brittle materials. In soil, geologic materials, and concrete, hardened steel
penetrators typically are used to penetrate in a non-deforming or rigid mode. The
benefit of this is that the penetration depth (typically normalized by length,
diameter, or cubed root of mass) is proportional to the ratio of the penetrator
strength to the target strength. Above the impact velocity where the penetrator
deforms (a measure of its dynamic yield strength), this penetration depth is
proportional to the square root of this ratio and approaches this ratio times the
penetrator length at high impact velocities, the often termed, hydrodynamic limit.
Another benefit to using these techniques is that the deformation and flow
characteristics of the penetrator do not affect the results. When the penetrator
deforms, this behavior must be understood to interpret the data obtained. Strain-
rate effects, adiabatic-shear failure, etc., all become necessary to quantify.
Traditional ballistic experimental methods: Typcially, when one thinks of
ballistic experimental methods, one thinks of four test techniques, three of which
are described below, the other is a ballistic limit test. Their description and
history could be, and perhaps should be, the entire focus of this paper. However,
we chose to provide a more complete listing and utilize extensive vuegraphs (a
subset appear in the Appendix) and references to steer the reader to the more
traditional experimental methods. The references provided will detail these
techniques in great detail, and most researchers and engineers within the armor
ballistics community are familiar with these techniques. The more recent
newcomer to this list includes penetrator dwell, which is a pre-penetration phase
that has very significant implications. We discuss these somewhat more
extensively later in this paper.
PEN: Semi-infinite Penetration: The generic ballistic material behavior of
a material is expressed in terms of a penetration vs. impact velocity curve. These
curves can be fit with empirical curves useful for systems-level modeling. Direct
and reverse ballistics techniques have been used extensively. Confinement is
necessary and sometimes affects the results. Nevertheless, a measurement of the
resistance to penetration, the consumption rate of the eroding penetrator, and of
the areal density per unit mass of the ceramic penetrated, are very useful to the
armor designer. An invited speaker details this technique, e.g., see Orphal, this
symposium.
Ceramic Armor Materials by Design 119
DOP: Depth-of-Penetration: In the mid-1980’s an attempt to standardize
an experimental method to rank ceramics for armor applications led to the
development of this technique. Utilized by numerous researchers, it is typically
the most-often used data, but the relevance of the data generated is highly
debated, basically because the configuration is not representative of most armor
applications. We present detail in our presentation, and an invited speaker will
discuss the issues associated with this common, but not standardized ballistic
testing technique. An international conference was held to discuss this technique
and attempted to standardize it, but generally each country uses its own
techniques, penetrators, target configuration, attention to interfaces, and whether a
cover plate is utilized or not. Reader is referred to James, this symposium.
DWE: Dwell: The discovery of dwell on the ceramic surface prior to
penetration is one of the most significant. Simply stated, the ceramic will resist
penetration, until it fails in some manner due to a number of possible reasons,
after which penetration will commence. George Hauver has gone through great
expense to carefully prevent the ceramic from failing, and was able to completely
erode any penetrator on the ceramic surface with no penetration whatsoever. This
is termed interface defeat, and would be the envy of all armor designers. A great
amount of research is being expended to understand this behavior and its
implications, which we believe are very significant. In fact, the techniques used
to interpret data obtained using traditional ballistic test techniques may need to be
re-examined, due to the fact that penetrator dwell is often present.
DPT: Dwell/Penetration Transition: As the impact velocity is increased, a
level is reached where the compressive strength of the ceramic is exceeded, and
dwell is no longer possible, or any significant duration of dwell. The
experimental technique used to obtain this information has been described in the
last three International Symposium on Ballistics and in the open literature. An
invited speaker will discuss this in more depth; see Lundberg, this symposium.
Armor Design-Oriented Ballistic Experiments
This category of experimental techniques represents armor system applications,
and hence it can be referred to direct evaluation of a material in the particular
application. Therefore, these are the most important results. However, these are
the most difficult test methods to utilize to provide information on how to
maximize the performance of the material. The probabilistic nature of these tests
also requires a large number of tests to be conducted. It is for this reason, that
simpler screening experiments described earlier have been developed. These
experiments are often used to validate numerical and analytic models, but this is
120 Ceramic Armor Materials by Design
cautioned, as discussed in Normandia at the 45th
US Army Sagamore Conference
on Armor Materials by Design.
Table III Categorization of Testing Techniques – Armor-Design Oriented
(Ballistic) Experiments and Methodologies
CATEGORIES OF CERAMIC TESTING UTILIZED TO EVALUATE
ARMOR PERFORMANCE OF CERAMICS
Armor Design-Oriented (Ballistic) Experiments and Methodologies
(8) (FTG) Fixed geometry (e.g., 1-4-3 thicknesses at 60-deg obliquity)
(9) (TCA) Tandem configurations (MTL/BRL patent)
(10) (VBL) Ballistic Limit Velocity tests (V50 or perforation test data)
(11) (BAD) Behind Armor Debris
(12) (TAD) Minimum Target Areal Density (different for each material combination)
a. (PAD) Protection Areal Density Testing
b. (TTE) Threshold Thickness Experiments
FTG: Fixed Target Geometry: Experimental techniques that attempt to
include the effects of finite thickness or impact obliquity are more representative
of the actual armor geometry and have been used to better compare and rank
ceramic materials for their intended use. Fixed geometry targets are utilized as a
method to rank materials in more-realistic armor environments, while
standardizing the test methods to avoid the probabilistic interpretations required
of the more traditional ballistic limit test methods. A common example is a target
denoted as 1-4-3, where the target consisted of a unit thickness metallic cover
plate, a ceramic of 4 times the unit thickness and a metallic backing plate of 3
times the unit thickness. Targets with this or other finite thicknesses (or weights)
were typically tested at normal impact, but some standardized testing has been
conducted at 60-degrees obliquity.
TCA: Tandem Composite Armor: The difficulty in utilizing ceramics in some
armor applications is the result of accumulative damage effects, and the pressure-
dependent behavior of comminuted and fragmented ceramic. Thus, thicker
ceramic tiles often do not perform as multiples of lesser thicknesses, making it
expensive and less mass efficient to use ceramics to defeat more lethal threats. A
technique to utilize a thinner ceramic tile backed by a metal backplate in a
repeated environment, was developed as a joint effort between the Ballistics
Research Laboratory and the Army Materials Research Laboratory, both now part
of the US Army Research Laboratory. The performance of the armor system
approached the multiplicative performance of each independent system, due to
isolation of each system from each other.
Ceramic Armor Materials by Design 121
VBL: Ballistic Limit Velocity: v50 or vL tests determine the effectiveness of the
armor system, often in the configuration actually used. Typically, a minimum of 5
experiments are required within a tight velocity range with at least 2 partial or 2
complete penetrations to determine a v50, and a much larger number of
experiments for a v95, typically used as acceptance criteria for actual deliverables.
BAD: Behind Armor Debris: Tests that determine the vulnerability of the
contents behind the armor evaluate the performance in an overmatched situation,
likely to occur in typical ballistic environments. Less debris is desirable for the
armor, more for the penetrator. Tests that quantify this effectiveness have been
developed and used for this purpose. The use of behind armor residual penetrator
and target debris is an often-used evaluation of lethality or vulnerability and is
critical in systems level evaluations.
TAD: Target Areal Density: Protection Areal Density (PAD) and Threshold
Thickness Experiments (TTE) have been developed to understand the general
behavior in ceramics to defeat a particular class of threat. These tests fix impact
conditions and adjust the target material allocation until a penetration threshold
target is obtained. This threshold target delineates the minimum weight for the
given thickness of ceramic that will defeat the threat for the particular material
allocation and impact conditions. Repeating this process for various ceramic
thicknesses generates a threshold curve. Understanding the general behavior of
this curve permits the designer to predict the performance against any threat, with
the caveat that the defeat mechanism has not changed as one changes the threat.
The threshold curve, which separates armor failures and successes (within some
defined probability), shows that there exists an optimum target configuration (a
particular material allocation) that minimizes the target areal density (weight) to
defeat the threat. The general behavior of targets near this optimum are examined
and used predict minimum target areal densities against various threats and
impact conditions. An invited keynote presentation that describes one of these
techniques will illustrate this procedure. Normandia has also presented an
alternative theory based on the relative time to failure of the confinement system
to that of the penetration time. The reader is referred to the keynote speaker,
Adams, this symposium, for discussion on the PAD technique.
While all of the test methods provide useful information, when taken together,
the results better define a material’s capability in a practical armor configuration.
The tests that are the more traditional ballistic experimental techniques appeared
in bold type in Tables II and III, and are re-tabulated in Table IV along with the
common test objectives and information provided. Most of these will be presented
in separate papers in this conference, as will some of the testing methods utilized
in the first two categories, as well as the traditional analytic and numerical tools
and methods used to evaluate the ballistic response.
122 Ceramic Armor Materials by Design
Table IV Ceramic Material Evaluation Summary of Ballistic Test Methods
TEST TEST TYPE INFORMATION OBTAINED REFNDP Non-Deforming
Penetration
Typically used for soft metals and hard targets, this
applies for concrete, limestone and other geological
materials. Various researches attempt to isolate
target resistance in this penetration mode.
10, 12,
49
PEN Penetration Depth
Direct or Reverse
Impact
Penetration-velocity curves, penetration resistance,
penetration rate, penetrator consumption rate.
21, 42-4,
57
DOP Modified Depth-of-
Penetration
Relevant for determination of performance goals as
a function of ceramic thickness – similar to TAD,
but in a semi-infinite configuration.
3, 4, 13,
18, 25,
39, 48,
50-1,
67-71
DWE Dwell Tests Total interface defeat conditions. 5-7, 9,
16-7, 19,
20,35-6,
41, 46-7,
60
DPT Dwell/Penetration
Transition
Velocity defines a load that is characteristic of a
failure shear strength of the ceramic, or of a
transition strain.
28-30,
32-34
FTG Fixed Target
Geometry
Generic material comparison experiment in armor-
like configurations, particularly at obliquity.
22-3, 72
TCA Tandem Composite
Armor
Configuration to minimize the use of damaged
material.
15,
VBL Ballistic Limit
Velocity (V50) or
Residual Data
Typical requirement for acceptable armor,
individual tests measure residual penetrator
characteristics.
14, 45,
55
BAD Behind Armor
Debris
Used to measure the lethality of the penetrator or
the vulnerability of the target to an overmatched
threat. Data quantification utilized in lethality
assessment tools.
TAD Target Areal
Density
Performance Maps
Helps determine near-optimal armor configurations.
Theories permit extrapolation to different threats.
Includes PAD, TTE, and TSM methodologies.
1, 8, 40
DISCUSSION
When dwell was first discovered, it was obtained under small-scale testing
in a well-controlled laboratory environment. Since that time, the dwell
phenomenon has both fascinated and captivated researcher’s interest over the last
decade, particularly when total interface defeat of a penetrator is achieved
[Hauver, et. al]. Tests that provide information about the dwell phenomenon have
Ceramic Armor Materials by Design 123
recently been added to the traditional suite of ballistic tests, and in our opinion,
are necessary to fully characterize the performance of ceramics for use in armor
applications, even if the dwell mechanism is not being utilized for threat defeat.
The successful up-scaling of interface defeat for bullet and medium-caliber
applications, albeit not in armor configurations, has led to a proliferation of
worldwide research activities. Subsequently, attempts to model this behavior
using the numerical tools and recent constitutive model formulations has also
accelerated, with limited success. The generation of data demonstrating dwell
and/or total interface defeat of bullets impacting ceramics and experiments to
determine the transition impact velocity between dwell and penetration for long-
rod, kinetic-energy penetrators has demonstrated that dwell is a natural occurring
phenomenon, which must be accounted for in material constitutive behavioral
models. The extensive data generated by Hauver, et al. over the past decade has
not yet been widely disseminated in the open literature and is being compiled into
an ARL report [Rapacki, et al].
Briefly, the possible occurrence of dwell is likely to have been present for
some duration, however slight, during most traditional ballistic testing techniques.
The lack of accurate accounting for the presence or duration of dwell during the
ballistic experiment has clouded the use of this data for constitutive, numerical or
analytic model validation. If dwell were present during the test, which we think
likely in most traditional tests, subtle geometric differences can result in variable
performance. For example, if dwell duration were to vary between tests due to
subtle geometric differences in target construction by several micro-seconds, one
would measure variable penetration resistances, which are likely to show up as
time-dependent or velocity-dependent target resistance. Generally terms that
account for the dwell phenomenon and its duration are lacking in the models used
to deduce target resistance. Without accounting for dwell, an exaggeration of the
variability of the measured target resistance is highly likely, and the performance
effects due to geometric differences are highly pronounced. An example of this is
attempting to compare ceramics using constant finite thickness (or constant areal
density) tiles. If ceramic bending were the cause of ceramic tensile failure,
different loads would produce this failure, and if this failure affected dwell
duration, ceramic comparative performance would be affected.
The reliance on numerical tools to guide armor development requires that
the material constitutive behavior be accurately described. The dependence on
the use of ballistic test results to determine model constants is disturbing.
Potentially worse, in our opinion, is the fact that dwell was likely present in some
the experimental data used to calibrate these model constants. Since the
numerical modeling of ceramics utilizes constitutive model formulations that may
not accommodate dwell, or the numerical methods do not allow dwell to occur for
any length of time, these constants may be determined incorrectly. Since the
124 Ceramic Armor Materials by Design
strength of the failed material was deduced from the experimental data that may
have had dwell, that value is also likely to be incorrect.
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Ceramic Armor Materials by Design 125
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126 Ceramic Armor Materials by Design
25 James, B., “The Influence of the Material Properties on Ballistic
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Ceramic Armor Materials by Design 127
37 Milman, Yu. V. and Chugunova, “Mechanical Properties, Indentation and
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128 Ceramic Armor Materials by Design
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Ceramic Armor Materials by Design 129
64 Wilkins, M. L., Honodel, C. A., and Sawle, D. R., “An Approach to the
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th Intl. Symposium on
Ballistics, Interlaken, Switzerland, 2001.
130 Ceramic Armor Materials by Design
Ballistic Test Techniques(IND) Indentation tests
• Micro-hardness indentations used by Milman to develop stress-strain relations, plasticity and dynamic hardness
• Hertzian indentation techniques used by NIST/Dr. Lawn andARL/Dr. Weresczak generates load/unload curve
• Used to screen and rank ceramics for armor
PushRod
DiamondHertzian
(spherical)Indenter
Specimen
CapacitanceDisplacementGages (3) &
Holder
Loads tohundreds of
pounds
Conducting Tape(i.e., 1/2 of parallel plate capacitor)
Ballistic Test Techniques(NDP) Non-Deforming Penetration tests
• Areal density penetrated DATA as a function of velocity usedsuccessfully to screen and rank armor ceramics
• Resistance to penetration is independent of penetrator strengthIF and WHEN penetration is rigid
• Theoretical expression estimated dynamic hardness• Rigid/deforming penetration models were used to compute
penetration resistance, which were about 1/3 of the theoreticalvalues computed (Sternberg, JAP 1989)
Metal backing
Ceramic
WCCeramicsphere
Heavily ConfinedMetal frame
Pen
Velocity
Vary velocityVary pressure
Analytic SolutionFor Rigid Penetration
Determines Target Resistance
DOP data
Ceramic Armor Materials by Design 131
132 Ceramic Armor Materials by Design
Ballistic Test Techniques(CEX) Cavity Expansion tests -SRI
• Expanding SRI cavity experiments funded by ARO in mid 90’s
• Collapsing cylinder experiments conducted by Nesterenko
• Cavity Expansion Theory computes the pressure at the penetrator/targetinterface from integration of static or dynamic stress fields in the target
• Linkage implies CEX tests can be used to directly measure penetrationresistance
HE
ceramictube
taperedbrass pin
Empty
space
foamattenuator
Collapsing tube experiment(variation of experiment by Nesterenko et al)
detonator
copperconfinementtubes
HE insphericalcavity
ceramic cylinder
brassconfinementcylinder
Expanding cavity experiment
Ballistic Test Techniques(PEN) Penetration tests: Direct Ballistics
• Direct penetration of thick, confined ceramics in either direct orreverse ballistic configuration
• Penetration time histories obtained from flash x-rays• Typical penetration vs. impact velocity curves obtained• Los Alamos National Laboratory Phermex Experiments
Ceramic Armor Materials by Design 133
Penetration TestsReverse Ballistics Configuration
D. L. Orphal and R. R. Franzen, “Penetration of Confined SiliconCarbide Targets by Tungsten Long Rods at Impact VelocitiesFrom 1.5 to 4.6 km/s,” Int. J. Impact Eng., v19 (1997) 1-13.
Tungsten (99.95%)ρ = 19,300kg/m3
L = 15.24mm or11.43mmD = 0.762mm
(PEN) “Long Rod Penetration of Ceramics,”
Dr. Dennis Orphal
Int’l Research Associates, CA
Ballistic Test Techniques(DOP) Depth-of-Penetration tests
• DOP technique developed to evaluate ceramic materials(ceramic effectiveness factors on an areal density basis)
• Data corrected for to nominal impact velocity of 1500 m/s• Initial tests: no cover plate, semi-infinite RHA backing,
bonded interface, W or DU, 65 g L/D 10 penetrators
L/D=1065 g
W or DU
“Depth of Penetration Testing,” Dr. Bryn James, DSTL, United Kingdom“Armor Ceramics Under High Velocity Impact of a Medium Caliber LRP,”Ernst, H., Wiesner, W., and Wolf, T., ISL, France
134 Ceramic Armor Materials by Design
Figure 2.1.6.1 Target andPenetrator Descriptions forSilicon Carbide DOPExperiments, Franzen et al.
Silicon Carbide
RHAbase target
Tungsten penetratorL = variedD = varied
Pr = residual penetration into RHA base target
tc
Figure 2.1.6.2 Target and PenetratorDescription for DOP Experiment,Rosenberg and Tsaliah [39].
Silicon Carbide(Mat. #118)
RHAbase target
Tungsten Alloypenetratorρ = 17800kg/m3
L = 80mmD = 8mm
12.7mm
Pr = residual penetration into RHA base target
Ceramic is bonded to theRHA base target
DOP Test Variations
SiliconCarbide(Mat. #113)
tc
HH RHAUTS = 1.45GPaRc~45
30mm
100mm
20m
m
10mmMild Steel
Tungsten Sinter-Alloy PenetratorV = 1700m/sρ = 17600kg/m3
UTS = 1.2 GPaelongation = 10%L = 72.5 mmD = 5.8 mmL/D = 12.5
1.5mm rubber
Pr= residualpenetrationinto HH RHA
Ceramic is gluedto RHA base
Target is square
Figure 2.1.6.4 Target and Penetrator Descriptionfor DOP Experiment, Rosenberg et al. .
Figure 2.1.6.3 Target and Penetrator Description for DOP Experiment, Reaugh et al.
Silicon Carbide (Mat. #110)
4340 Steel Base TargetRc = 33-37
Tungsten Sintered Alloy W2 Penetratorρ = 18360kg/m3 Rc = 28-31L = 25.4mm UTS = 0.88GPaD = 6.35mm Yield = 0.695GPaL/D = 4 Elongation to fracture = 5.5%
tc
Pr = residual penetration into RHA base target
Ceramic is bonded to theRHA base target using Stycast 1266
102 mm
152mm
64mm
V
T. Holmquist, R. Rajendran, D. Templeton, &, K. Bishnoi,“A Ceramic Armor Materials Database,” TARDEC report13754, Warren, MI, Jan. 1999
• George Hauver discovery began by determining the change intarget strength as a function of time
• Achieved interface defeat under strong confinement• Extensive recent activity has significant implications
Ballistic Test Techniques(DWE) Complete Dwell tests
TiC
Ceramic Armor Materials by Design 135
Dynamic Confinement TestsSabot Preloads Ceramic
Tranchet and Malaise(Centre d’Etudes de Gramat and ENSAM/LAMEF, Talence, France)CMWG mtg. 8-10, 1999 (courtesy of Dr. Bless)
Ballistic Test Techniques(DPT) Dwell/Penetration Transition tests
• Similar to other reverse ballistic experiments in design,however, controlled heat-shrunk confinement and frontplate shock-mitigation was used
• Penetration-time histories established dwell occurrence andtransition to penetration
Rear Plug
Front Plug
Locking Rings
Tube
12φ
20 8
Ceramic
2028
Rear Plug
Front Plug
Locking Rings
Tube
12φ
20 8
Ceramic
2028Dimensions in millimetersTempered steel front and rear plugs
(750 MPa flow stress).Maraging steel tube and locking rings
(Mar 350, 2600 MPa flow stress).
P. Lundberg, R. Renstrom, & B. Lundberg, “Impact of Metallic Projectiles onCeramic Targets: Transition Between Interface Defeat and Penetration,” Int. J.Impact Eng., v24 (2000) 259-275.
P. Lundberg, R. Renstrom, & L. Holmberg, “An Experimental Investigation onInterface Defeat at Extended Interaction Time,” Proceedings of the 19th
International Symposium on Ballistics, edited by I. Rose Crewther, v3, 1463-1469.
(DWE) “An Analysis of the Transition between Interface Defeat and Penetration forA Given Combination of Projectile and Ceramic Material,” P. Lundberg, R.Renstrom, L. Westerling, Swedish Defense Research Agency, Sweden
136 Ceramic Armor Materials by Design
Ballistic Test Techniques(FTG) Fixed Target Geometry tests
• ARL 1-4-3 tests in 1980’s– Determined standard limit velocity VL
– Equation estimates zero residual penetration velocity– 3 or 4 tests for each ceramic at normal obliquity
• EMI/DERA tests at obliquity– Determined optimum material allocation and weights
Ceramic25.4 mm(nominal)
Metal backing19-mm
High Hard Steel(or ESR or RHA)
Metal cover6.25 mm RHA
19-mmMild Steel
Confinement
Metal backing
Ceramic Tilesat obliquity
Penetrator
V. Hohler, K. Weber, R. Tham, B. James, A.Barker, &, I. Pickup, “Comparative Analysis ofOblique Impact on Composite Systems,” Int. J.Impact Eng., to appear, HVIS 2000proceedings.
ARL Standardized Ceramic Target: ATM-C
Ceramic Armor Materials by Design 137
AD995 Ceramic Disk101.6mm diameter
12.7mm thick
Ti-6Al-4V Cover Plate127mm diameter
6.35mm thick
Ti-6Al-4V Confinement Ring
Ti-6Al-4V Cover Plate127mm diameter
6.35mm thick
Standardized Research TargetInitial Encapsulated Ceramic Configuration
Electron-beam weldedin vacuum
Representative,IdealizedConfiguration
ARL/NIST/SNL/LANL effort
Ballistic Test Techniques(TCA) Tandem Composite Armor tests
• BRL/MTL Damage Mitigation Configuration• Design approach isolates ceramic material to repeatedly achieve
material performance potential– ad hoc, intelligent solution that capitalized on empirical
testing observations– Trades off space for performance– Good performance levels achieved
• Component nature of design provides design flexibility
Metal backing
Ceramic honeycomb
Test bed for materials, models, and design optimization
138 Ceramic Armor Materials by Design
Ballistic Test Techniques(VBL) Ballistic limit velocity tests
• Probably the most extensively used data technique• Statistical data analysis required multiple tests• Higher confidence levels requires significantly greater number
of experiments• Used for acceptance testing and is standardized• AMTL Mascianica Handbook available in digitized form
Silicon Carbide(Mat. #105)
6061-T6 Aluminum Back Plate
6.35mm
6.35mm
Allegheny Steel 609 Sharp PenetratorRc = 54-56M = 8.32gL ~ 29 mmD = 7.62 mmcone angle = 55o
Target Configuration
Figure 2.1.7.1 Target and Penetrator Description forPerforation Experiment, Wilkins et al. [26].
Bonded using polyurethane(Scotchcast 221)
Ballistic Test Techniques(TAD) Target Areal Density tests
• Systematic ballistictesting was used todevelop armor designmethodologies during the1980’s– JPL (figure shown)– DARPA contractors
• Honeywell (Alliant)• DuPont and GD teams• A.R.A.P./Abex/Norton
– LANL/LLNL
(TAD) “Theory and Experimental Test Methodsfor Evaluating Ceramic Armor Components,”Dr. Marc Adams, Jet Propulsion Laboratory, CA
THEORY AND EXPERIMENTAL TEST METHODS FOR EVALUATING
CERAMIC ARMOR COMPONENTS
Marc A. Adams
Jet Propulsion Laboratory
4800 Oak Grove Drive, Bldg 122-B3
Pasadena, CA 91109
ABSTRACT
The Ballistic Performance Map (BPM) and its derivatives are useful
constructs for understanding the function and performance of ceramic and other
components in hardfaced armor systems. By combining the BPM model with a
constant velocity ballistic testing approach and varying the areal densities of the
target components, the Protection Areal Density (PAD) testing methodology, the
relative and absolute performance of different ceramic materials are readily
evaluated. In addition, such methods enable adequate statistical analysis of the
unbiased test results to understand the basic uncertainty in the measured
performance and the stochastic behavior of any hardfaced armor material or
armor system. Examples of the ballistic performance of several classes of
hardface materials are given.
INTRODUCTION
This paper discusses the results of studies, sponsored by the U.S. Army
TACOM, that were conducted in the late 1980's and early 1990's. These studies
developed new methods for the ballistic evaluation of candidate armor materials
and evaluated the ballistic performance of a variety of ceramic materials. Before
initiating the experimental evaluation program, the various ballistic test methods
being used at the time for evaluation of armor materials and components were
critically examined. None were found to be adequate for the statistically
meaningful characterization of ceramic material performance. Often, large
uncertainties in the measured performance values were ignored and bias in the
performance measures, introduced by the testing methods, was not being
adequately addressed.
Ceramic Armor Materials by Design 139
To the extent authorized under the laws of the United States of America, all copyright interests in this publication are the propertyof The American Ceramic Society. Any duplication, reproduction, or republication of this publication or any part thereof, withoutthe express written consent of The American Ceramic Society or fee paid to the Copyright Clearance Center, is prohibited.
A new method for the ballistic evaluation of materials, components and armor
systems was developed. In addition to developing the experimental procedures
and statistical experiment design philosophies, analysis methods were developed
that enable the statistical interpretation of experimental test results and
identification of uncertainties associated with the measured values. This testing
method is called PAD, the determination of the protection areal density of a
material, component or armor system. The method fixes the test projectile
characteristics, e.g. projectile type, impact velocity, obliquity, yaw, and varies the
areal density of the targets used in the test series. The fundamental determination
made from the ballistic test data is the probability of partial or complete
penetration of a target of given design and areal density. The data is analyzed
with binomial statistical procedures or by maximum likelihood estimation
techniques.
The ceramic material evaluation studies characterized the performance of
ceramic and ceramic composite materials for use as components in armor systems
to protect against "small arms, kinetic energy threats". Armor piercing (AP)
projectiles with hard penetrator cores were used as the test projectiles. The .50cal
AP M2 projectile was used for the majority of the studies, although some
evaluations were also performed with .30cal and 14.5 mm AP projectiles. Over a
period of several years, ceramics companies, ceramic developers at universities
and government laboratories submitted specimens of various types of ceramic and
ceramic composite materials that were evaluated using the methods described in
this paper. The results of some of these evaluations are presented below.
USE OF CERAMICS IN ARMOR SYSTEMS
Ceramic containing armor systems typically have configurations similar to
that shown in Figure 1. The ceramic or cermet (hardface) component is usually
one of the first armor system components impacted by the projectile. Situated
behind the hardface component are one or more backing components that provide
support to the brittle hardface plate and affect the final defeat of the damaged
projectile and the ceramic debris. The shape and dimensions of the hardface
material vary from one armor system to another but the basic function of the
hardface remains that of damaging (cracking, shattering, eroding) the incident
projectile and turning or yawing the projectile from its incident trajectory.
Other features are often incorporated into ceramic armor systems. Typically,
a spall shield component is placed in front of the ceramic to suppress the ceramic
debris thrown off the front face during projectile impact. The ceramic is attached
to the backing component with adhesive or by other means. Many armor systems
have a requirement to defeat multiple hits of the threat projectile, some hits in
close proximity. Often, the individual plates of ceramic are structurally isolated
140 Ceramic Armor Materials by Design
with material in the area between the plates to prevent adjacent ceramic plates
from being damaged by a hit. Additional dynamic isolation may be required
Ceramic Plates
Metallic or Polymer Composite Backing Plate(s)
Bond of Ceramic to Backing Plate
Material isolates ceramic plates
Thin Cover Plate
Figure 1. Typical Configuration of a Ceramic Armor System
between the ceramic plate and the backing components. Alternately, "tough"
ceramics and cermets have been investigated for their ability to defeat impacts of
projectiles and limit the lateral damage created in the ceramic such that
subsequent, proximate hits of the threat can be defeated by a single continuous
ceramic plate instead of the "tiled" array of plates shown in Figure 1.
UNDERSTANDING THE PERFORMANCE OF CERAMICS IN ARMOR
SYSTEMS
In general, the performance of a hardface material in an armor system cannot
be predicted from the intrinsic properties of the hardface material at the present
time. Few static properties correlate with the ballistic performance of ceramic
materials. Comprehensive physical models of the penetration event, which use
intrinsic material properties and closed form physical descriptions, are not capable
of accurately predicting the marginal conditions under which a given projectile
will completely penetrate the armor. Measurement of relevant dynamic material
properties sheds some light on the suitability of a ceramic material but the
relationship of these properties to actual performance in an engineered armor
system cannot be relied upon, at present, for design purposes. Some success has
been obtained in efforts to use discretized, finite element/finite difference
modeling or "hydrocode" modeling but these analytical tools are not adequate to
design armor systems or to adequately predict material performance. They simply
don't describe, with sufficient accuracy, all of the important phenomena that affect
performance. The basic dynamic material behavior, under the conditions of
projectile impact and penetration, is still imperfectly understood.
In order to accurately characterize the relative or absolute performance of
ceramic materials for use in armor systems, it remains necessary to evaluate the
materials by ballistic testing of the armor system or testing a target that is a good
Ceramic Armor Materials by Design 141
surrogate of the system. All important phenomena that occur in the armor system
must occur during the projectile penetration of the target used for the evaluation;
the target design must insure this. Such testing is more akin to component
evaluation than to the determination of intrinsic material properties. The design of
the target influences the absolute performance level of the ceramic component
and affects any performance comparisons with targets using other materials.
EVALUATION OF THE BALLISTIC PERFORMANCE OF CERAMICS
All of the design features described above affect the ballistic performance of
the hardface armor component. Depending on the particular design of the armor
system, the hardface component may be more or less effective in damaging the
projectile and contributing to its defeat. The development of any efficient armor
system requires the complex, co-optimization of several design parameters and
materials selections. This fact complicates attempts to evaluate the relative
performance of ceramic materials tested in different target designs and
complicates evaluation for different armor system designs. A fixed target design
and test projectile should be used for comparative ballistic evaluation of hardface
materials. This target design must faithfully create the same penetration
phenomena as the armor system for which the ceramics are being evaluated.
Given that ceramic materials are best evaluated as "components" in a system,
the most unambiguous performance measure is whether the target used for the
evaluation, is partially penetrated or completely penetrated by the projectile.
Binomial data is taken in such tests. The target of a given configuration is either
partially penetrated or completely penetrated by a given projectile impacting at a
given velocity and obliquity. Other diagnostics can be used in these tests such as
capture of the damaged projectile to determine the level of breakup and
measurement of the deformation produced in the backing plate of partially
penetrated targets.
Ballistic testing approaches typically vary one of two principal variables in the
course of the experimental determination. The target design can be fixed and the
impacting velocity of the projectile varied. This testing is called ballistic limit
determination or the determination of "V50", by definition the velocity at which
there is a 0.5 probability that the target will be completely penetrated. The other
approach is to fix the impact velocity of the chosen test projectile and vary the
areal density (thickness) of the target components. The probability of partial
penetration as a function of target areal density is determined. This is the PAD
method of ballistic performance characterization.
The requirements for most new armor applications define the threat
projectile(s) and impact velocity(s). The program goals are usually to find the
lightest weight armor system; minimizing armor system areal density is of
greatest interest. Curiously, much of the ballistic testing conducted in these
142 Ceramic Armor Materials by Design
programs employ the variable velocity approach and determine V50 quantities.
The relationships between target design, target areal density, impact velocity and
defeat of the projectile are complex. It is best to measure, directly, the principal
variable of interest and hold all other variables as constant as possible.
Figure 2 illustrates the nature of the two different test methods and the
features of the Ballistic Performance Map. On the base of the three axes plot are
the areal densities of ceramic and backing in the target. The vertical axis
represents the impact velocity of the projectile. The ballistic limit surface, shown
with constant velocity contours, represents the locus of Target Design Points
(ceramic and backing areal density) that will defeat the projectile (prevent
complete penetration of the target) at the given impact velocity 50% of the time.
Target tests used for the PAD test method and for the variable velocity test
method are shown in the cutaway area. The stochastic behavior of the complete
penetration event can be visualized as a "thickness" of the ballistic limit surface.
As one moves upward through the surface from below, the probability of
complete target penetration increases from near zero toward unity. The Ballistic
Performance Map is the base of this figure. The PAD line for an impact velocity
is the projection of that constant velocity contour of the ballistic limit surface onto
the base of the figure.
Figure 3 illustrates the design of PAD tests and the methods used to analyze
the ballistic test data. Appropriate Test Line(s) are chosen and the test lines are
populated with targets at selected Target Design Points. The density of Target
Design Points on the Test Line is a balance between the number of targets that
can be used and the degree of accuracy and statistical confidence required for the
determination of the partial penetration probability along the Test Line. Binomial
statistical analysis can be used to establish the uncertainty interval for the
measured value of partial penetration probability at each target design point.
Alternately, maximum likelihood estimation techniques can be used to analyze
the data on the Test Line to obtain the partial penetration probability (with its
uncertainty) as a function of target areal density on the Test Line. These test and
analysis methods were used to obtain the results presented in this paper.
Ceramic Armor Materials by Design 143
Figure 2. Ballistic Limit Surface and Ballistic Performance Map
144 Ceramic Armor Materials by Design
THE BALLISTIC PERFORMANCE OF VARIOUS CERAMIC MATERIALS
The target chosen for the ceramic evaluation program was a 4x4 inch plate of
ceramic bonded with solid film adhesive to an octagonal shaped backing plate
made from 5083- H131 aluminum alloy. If the ceramic test plate was not flat to
within 0.003 (in), the minimum required thickness of polyester resin was cast onto
the surface to make it flat and eliminate the possibility of voids in the bond line.
Material suppliers were requested to provide each of their ceramic plates in a
thickness that gave each plate an areal density of 11 (lb/ft2). The testing was
performed along a constant ceramic areal density Test Line. The backing areal
density was varied over a range that produced partial and complete penetrations in
the targets. The tests of the highest areal density targets with complete
penetration and the tests of the lowest areal density targets with partial penetration
were replicated as many times as possible to increase the statistical resolution and
decrease the uncertainty in the PAD determination. Measurements of the
permanent deformation of the aluminum backing plates were made for all
partially penetrated targets. Also, the degree of projectile core breakup was
measure for all tests in which the front and rear target containment boxes
adequately captured the pieces of the damaged projectile.
Figure 3. Design and Analysis of PAD Tests
Ceramic Armor Materials by Design 145
Maximum likelihood estimation (MLE) techniques were used to analyze the
data set of partial and complete penetrations obtained for each material. MLE
was used to determine: (i) the expected value of target areal density for a
probability of partial penetration (Pp) equal to 0.5 and (ii) the 90% confidence
interval of areal density for Pp = 0.5. Some materials have a large uncertainty
(confidence interval) associated with their expected PAD0.5 values. This arose
from two sources. In some cases, insufficient material was supplied to perform an
adequate number of tests to reduce the statistical uncertainty. In other cases,
sufficient material was supplied to test an adequate number of specimens but the
material's behavior was not consistent. These two cases cannot be distinguished
in the data presented. In general, poorer performing materials (ones that required
higher target areal density to defeat the projectile) also demonstrated more
inconsistency in their performance and had larger uncertainty bands.
In all, thirty-six materials were evaluated including: aluminum oxide, boron
carbide, boron carbide/aluminum cermet, silicon carbide and aluminum nitride.
The materials were provided by nine different organizations. As indicated on the
figures presenting the analyzed data, both hot pressed (hp) and sintered (s)
materials were provided.
Figures 4 through 7 summarize the results of the experimental investigations.
These results are plotted on Ceramic Performance Maps that display target areal
density as a function of ceramic areal density. Ceramic materials that have their
PAD points at lower target areal densities are higher performing ceramics.
Figure 4 presents the results for the five sintered aluminum oxide ceramics
evaluated at approximately 11 (lb/ft2). The "AD995 Al2O3 (s)" material is Coors
sintered, CAP3 alumina that was used extensively in the program to study the
effects of many variables on the performance of hardface components in armor
systems. It is one of the two baseline materials for the study and its performance
has been characterized in every area of the Ballistic Performance Map for four
different velocities. In each plot summarizing the performance data for a
particular class of material, there is a heavy orange line represented the expected
PAD0.5 value for the AD995 Al2O3 (s) ceramic and two thinner orange lines
representing the upper and lower bounds of the 90% confidence interval for the
PAD0.5 value. Similar lines are presented in blue on each plot for the other
baseline material, a hot pressed boron carbide made by the Dow Chemical
Company.
The aluminas [H] and [V] have performance indistinguishable from the
baseline alumina. Material [B] most likely has a PAD0.5 slightly less than the
baseline alumina. Material [A], a lower grade alumina, demonstrated inferior
performance to the other aluminas. Even this compositionally inferior alumina
requires less than 10% higher areal density to equal the performance of the better
aluminas.
146 Ceramic Armor Materials by Design
Figure 4. Expected Value and 90% Confidence Interval for
the PAD (Pp = 0.5) Values for Aluminum Oxide Ceramics
[B]
Al 2
O3(s
) [V]
Al 2
O3(s
)
[H]
Al 2
O3(s
)
[A] Al2O3(s)
AD995 Al2O3(s)
Baseline
80% confidence
AD 995
Al2O3
Dow B4C
15
16
17
18
19
20
21
10.7 10.8 10.9 11 11.1 11.2
Ceramic Areal Density (lb/ft^2)
Tar
get
Are
al D
ensi
ty (
lb/f
t^2)
Figure 5 summarizes the results of the evaluation of boron carbide containing
materials, two pure hot pressed boron carbides and one boron carbide/aluminum
cermet material. The two hot pressed materials have indistinguishable
performance that is about two lb/ft2 less than the baseline alumina. This
represents the performance difference between the best composition ballistic
ceramic for AP threats, B4C, and the lowest performing composition of ballistic
ceramic material, Al2O3. Several varieties of the B4C/Al cermet material were
evaluated. The one shown in Figure 5 was the best performer and, importantly,
was the most "ceramic like" in composition and microstructure of all the varieties
evaluated. It was found that, as the ductile metallic phases present in the cermet
body were increased, the toughness and lateral damage resistance increased;
however, the areal density required to defeat the threat, PAD0.5, increased
dramatically.
Figure 6 summarizes the results of the evaluation of six sintered and two hot
pressed aluminum nitride ceramics. No performance difference was observed
Ceramic Armor Materials by Design 147
Figure 5. Expected Value and 90% Confidence Interval
of PAD (Pp = 0.5) Values for Boron Carbide Ceramics
and Cermet
Dow
B4C
(hp)
[M] B4C/Alcermet
[S]
B4C
(h
p)
80% confidence
AD 995
Al2O3
Dow B4C (hp)
13
14
15
16
17
18
19
10.2 10.4 10.6 10.8 11 11.2
Ceramic Areal Density (lb/ft^2)
Tar
get
Are
al D
ensi
ty (
lb/f
t^2)
between the hot press and sintered materials. All materials have PAD0.5 values 1
to 1.5 lb/ft2 lower than the baseline alumina. One hot pressed material, [Q],
fabricated with "improved" techniques, had such inconsistent performance that
statistical analysis could say little about its performance other than it was
extremely inconsistent. Several of the sintered materials represent compositions,
microstructures and processing that were painstakingly developed over a
considerable period of time to be "superior ballistic materials". All this
development was based on static property measurements. None of these materials
are better than [C], the cheapest material with the simplest processing and a larger
grain size. It is not easy to change the composition, microstructure or processing
of a ceramic body and improve its ballistic performance. Ballistic testing during
material development is absolutely required to guide the development.
Figure 7 summarizes the results of the evaluation of six sintered and three hot
pressed silicon carbide ceramics. The performance differences between all but
one of the materials are small. Inspection of the expected values of PAD0.5, the
148 Ceramic Armor Materials by Design
Figure 6. Expected Value and 90% Confidence Interval for the PAD
(Pp = 0.5) Values of Various Aluminum Nitride Ceramics
[F]
AlN
(s)
[X]
AlN
(s)
[P]
AlN
(s)
[N]
AlN
(s)
[Y] AlN(s) 70%
confidence
80% confidence
[L] AlN(hp)
AD 995 Al2O3
Dow B4C
[C]
AlN
(s)
80% confidence
[Q]
AlN
(hp)
14
15
16
17
18
10.6 11 11.4 11.8 12.2
Ceramic Areal Density (lb/ft^2)
Tar
get
Are
al D
ensi
ty (
lb/f
t^2)
green data points, shows a general drift to higher target areal densities with
increasing areal density of the ceramic. This is an effect of target design and
points out the problem of comparing ceramic performance between ceramics
Ceramic Armor Materials by Design 149
Figure 7. Expected Value and 90% Confidence Interval for the
PAD (Pp = 0.5) Values for Various Silicon Carbide Ceramics
[O]
SiC
(s) [W
] S
iC(s
)
[G]
SiC
(hp)
80% confidence
[D]
SiC
(s)
[K]
SiC
(hp)
70% confidence
[R] SiC(s)
AD 995
Al2O3
Dow B4C
[I]
SiC
(s) [U
] S
iC(h
p)
[Z]
SiC
(s)
70% Confidence
14
15
16
17
18
19
20
21
10 10.5 11 11.5 12 12.5 13 13.5
Ceramic Areal Density (lb/ft^2)
Tar
get
Are
al D
ensi
ty (
lb/f
t^2)
evaluated in different designs. As the areal density of the ceramic is increased in
the target, the areal density of the backing decreases to maintain a constant target
areal density. The weight fraction of ceramic in the target increases, reducing the
weight efficiency of the target, irrespective of the ceramic performance.
Translation of the blue baseline performance line up or the orange baseline down
shows that the increase in target areal density is due to the target design factor and
not the fact that the higher areal density ceramics have lower performance.
150 Ceramic Armor Materials by Design
LONG ROD PENETRATION OF CERAMICS
D. L. Orphal
International Research Associates
4450 Black Ave.
Pleasanton, CA 94566
ABSTRACT
Reverse ballistic experiments were performed to measure penetration of long
tungsten rods into confined boron carbide, silicon carbide, and aluminum nitride
targets. Penetration depth and the length of the remaining rod were measured as
functions of time using flash X-rays. These data were used to determine the
velocity of penetration and the rate of rod erosion. Impact velocities ranged from
about 1.5 to 5 km/s. The experiments were performed using a two-stage light-gas
gun.
INTRODUCTION
Hohler and Stilp [1,2], Sorensen, et al. [3], and others have published data for
the penetration of long rods into steel and aluminum as a function of impact
velocity. The principle objective of this paper is to present similar data for
penetration of long tungsten rods into three confined ceramics; boron carbide
(B4C), silicon carbide (SiC) and aluminum nitride (AlN). The experiments were
performed in the reverse ballistic mode with multiple flash X-rays of the
penetration process. This approach results in data for both penetration depth and
length of eroded rod as a function of time. These data are used to determine the
velocity of penetration and rate of rod erosion, in addition to final penetration
depth. Impact velocities ranged from about 1.5 to 5 km/s.
A secondary objective is to briefly discuss the advantages and disadvantages of
reverse ballistic testing. Reverse ballistic testing is not a new technique but is
probably not used as often as would be beneficial.
EXPERIMENT DESIGN
These ceramic penetration experiments were designed to achieve several
objectives. Impact velocities were to range from about 1.5 to 5 km/s. Penetration
depth, p, and remaining rod length, Lr, were to be measured as a function of time,
Ceramic Armor Materials by Design 151
To the extent authorized under the laws of the United States of America, all copyright interests in this publication are the propertyof The American Ceramic Society. Any duplication, reproduction, or republication of this publication or any part thereof, withoutthe express written consent of The American Ceramic Society or fee paid to the Copyright Clearance Center, is prohibited.
allowing calculation of the velocity of penetration, u = dp/dt and the
“consumption velocity, vc = dLc/dt, where Lc = “consumed rod length” = L-Lr and
L = original rod length.
Given these objectives, it was decided to conduct the experiments in the
reverse ballistic mode. In direct ballistic experiments, the projectile, here a long
tungsten rod, is launched and impacts a stationary target. In reverse ballistics the
“target”, here a confined ceramic, is launched and impacts a stationary projectile.
In reverse ballistic experiments the size of the target is restricted by the size of the
gun and is thus necessarily small. The size of the projectile must be
correspondingly small. Thus reverse ballistics experiments are nearly always
small scale and the issue of scaling is very important.
A big advantage of small-scale reverse ballistics experiments is that, properly
designed, flash X-rays can be used to view the penetrator inside the target during
the penetration process. Multiple, independently timed flash X-rays provide
measurements of penetration depth and remaining penetrator length at known
times. In addition, the X-rays provide data relevant to the overall phenomenology
of the penetrator-target interaction, target hole size and growth, and the spatial
distribution of the eroded penetrator material. All these data are very difficult to
obtain in large-scale direct ballistics tests where X-rays cannot penetrate the target
and data is usually limited to final penetration depth and target hole geometry. An
excellent recent example of the use of reverse ballistic testing to study details of
the penetration process, specifically interface defeat by ceramics, is the work by
Lundberg, et al. [4] and Westerling, et al. [5].
Reverse ballistic testing has several other advantages. Penetration by very
complex penetrators, which would be very difficult to launch directly, can be
studied. Also, in reverse ballistic testing parameters such as angle of attack can be
precisely controlled. In the experiments reported here angle of attack was zero.
Target Geometry
A disadvantage of reverse ballistic testing is that because of the small scale and
the requirement to launch the target, target complexity is limited. In these tests
this limitation was not an issue. The ceramic targets were simple cylinders. The
diameter of the ceramic was selected to be as large as possible within the
limitations of the diameter and launch mass capability of the two-stage light-gas
gun. The targets for each of the three ceramics were basically the same, but
dimensions varied somewhat because of the different ceramic densities. Details of
the targets for each of the ceramics are reported in [6-8]. Here the B4C targets
used for v < 4.2 km/s are shown in Figure 1 to illustrate the essential features of
the target design. The longer target was used for lower velocity tests (1.5 v
2.77 km/s). The targets are composed of a ceramic cylinder radially confined by a
thin titanium sleeve and with 6061-T6 aluminum front and rear plates. Each
152 Ceramic Armor Materials by Design
titanium sleeve was machined to achieve a tight press fit with its matching
ceramic cylinder. The targets were surrounded by a 38 mm. outer diameter Lexan
sabot. In all tests the tungsten rod was completely eroded and the hole in the
target fully formed well within the ceramic.
Fig. 1. Typical target geometry (dimensions in inches).
Ceramics
The B4C ( t = 2.51 g/cm3) was hot-pressed by the Norton Company with a
typical grain size of 9 m. The SiC ( t = 3.22 g/cm3) was “pressure assisted
densified” (PAD) by Cercom; typical grain size was about 2 m. The AlN ( t =
3.25 g/cm3) was hot-pressed by the DOW Company and had a typical grain size
of 1.5 m. Additional information on the ceramics tested is given in [6-8].
Penetrators
Penetrators were long rods (right circular cylinders) of pure tungsten ( p = 19.3
g/cm3). Penetrator diameter, D, was selected to be visible in the flash X-rays and
small enough to insure a sufficiently large ratio of ceramic to penetrator diameter.
Littlefield, et al. [9] and Anderson, et al. [10] performed numerical simulations for
L/D = 20 tungsten alloy penetrators symmetrically impacting armor steel targets
that show no significant effects of the radial boundaries if the target diameter is
greater than about 15 to 20 penetrator diameters.
Ceramic Armor Materials by Design 153
For nearly all the tests reported here the rod diameter was 0.762 mm (0.030
inch). For some of the B4C tests in the velocity range 1.493 < v < 2.767 km/s a
penetrator diameter of 1.02 mm (0.040 inch) was also used. Thus the ratio of
ceramic to penetrator diameter was about 30 for nearly all the tests and was never
less that 20. In addition, examination of the flash X-rays from the tests reveals no
measurable radial expansion of the targets. Therefore it is believed that the targets
behaved as “well confined” targets.
Instrumentation
Primary instrumentation was four independently timed 450kV flash X-rays that
viewed the rod-target interaction. In addition, two continuous X-rays positioned
up-range of the impact and separated by 0.30 m were used to determine impact
velocity. The four flash X-rays contained fixed spatial fiducials and this plus the
known times the x-rays fired provided an independent measure of impact velocity
as well as the absolute zero time of impact. Impact velocities determined by the
independent measurements were always in excellent agreement.
ANALYSIS OF FLASH X-RAYS
Fig. 2 is a typical set of flash X-rays (Test 117, B4C at 3.134 km/s). The depth
of penetration, p, and length of remaining rod, Lr, is measured in each X-ray. The
length of rod “consumed, Lc, is Lc = L - Lr . Measured values for p and Lc for
Test 117 are given in Fig. 3. These data include the (0,0) point since time of
impact is independently measured.
Fig. 2. Typical flash X-rays Fig. 3. Data from flash X-rays
154 Ceramic Armor Materials by Design
The data shown in Fig. 3 are used to determine the following parameters:
Penetration Velocity, u
As can be seen from Fig. 3 the first four points in the penetration depth-time
plot, including (0,0), are well represented by a straight line of slope dp/dt = 2.039
km/s = u, with a correlation coefficient, r = 0.999. This slope is defined as the
penetration velocity.
Consumption Velocity, vc
The first four points on the Lc versus time plot, including (0,0), are also well fit
by a straight line of slope dLc/dt = 1.111 km/s = vc, with r = 0.999. This slope is
defined as the “consumption velocity” for the rod.
Primary Penetration Depth, pprimary
To a good first approximation it may be assumed for these high L/D rods that
the rear of the rod does not significantly decelerate until it reaches the target
interface. With this approximation, and assuming u and vc are constant as shown
by the data, the rear of the rod reaches the target interface at time tc = L/vc. The
depth of penetration at time tc is defined as the primary penetration, pprimary = utc,
and is shown as the open circle in Fig. 3.
Total Penetration Depth, ptotal
For each test one of the flash X-rays was timed to fire long after the penetrator
was fully consumed. The depth of penetration in this X-ray is defined as the total
penetration depth, ptotal.
EXPERIMENTAL DATA AND ANALYSIS
Due to space limitations the general phenomenology observed in the
experiments is not discussed but can be found in [6-8]. In discussing the
experimental data comparisons will be made to “ideal hydrodynamic theory” [11]
for which: uhyro = v / [1+ ( t/ p)1/2
] and phydro = L ( p/ t)1/2
.
Penetration Velocity, u
Figures 4, 5, and 6 show u versus impact velocity for B4C, SiC and AlN,
respectively. For these ceramics it is observed that penetration is steady-state (i.e.
u = constant) to a high degree of approximation. For the lowest impact velocities
u sometimes appears to be slightly non-constant and the lowest correlation
coefficients for a linear fit to the penetration-time data typically occurs for the
lower impact velocities. For the data shown here r 0.990 for a linear fit to the
penetration-time data in all cases.
Ceramic Armor Materials by Design 155
The dashed line is uhydro. As shown, for these ceramics u < uhyro over the entire
range of impact velocity. For each of the ceramics u is well fit by a linear
equation in v:
B4C: u = 0.757v - 0.406 (km/s) (1)
SiC: u = 0.781v - 0.510 (km/s) (2)
AlN: u = 0.792v - 0.524 (km/s) (3)
0.0
1.0
2.0
3.0
4.0
1.0 2.0 3.0 4.0 5.0
Impact Velocity, km/s
Pen
etr
ati
on
Velo
cit
y, u
,
km
/s
Fig. 4. Penetration velocity versus impact velocity for B4C.
0.0
1.0
2.0
3.0
4.0
1.0 2.0 3.0 4.0 5.0
Impact Velocity, km/s
Pen
etr
ati
on
Velo
cit
y, u
,
km
/s
Fig. 5. Penetration velocity versus impact velocity for SiC.
156 Ceramic Armor Materials by Design
0.0
1.0
2.0
3.0
4.0
1.0 2.0 3.0 4.0 5.0
Impact Velocity, km/s
Pen
etr
ati
on
Velo
cit
y, u
,
km
/s
Fig. 6. Penetration velocity versus impact velocity for AlN.
Consumption Velocity, vc
Plots of vc versus v are given in [6-8] but are not included here because of
space limitations. For each of the ceramics vc > vchydro = v - uhydro over the entire
range of impact velocity. There is significant scatter in the vc data but for each
ceramic the data are well fit with a linear equation in v:
B4C: vc = 0.219v + 0.333 (km/s) (4)
SiC: vc = 0.240v + 0.383 (km/s) (5)
AlN: vc = 0.216v + 0.434 (km/s) (6)
Primary Penetration, pprimary
Figures 7, 8, and 9 show pprimary normalized by original rod length versus
impact velocity for the three ceramics. The horizontal dashed line on each plot is
the ideal hydrodynamic penetration, ( p/ t)1/2
. For each of the ceramics pprimary/L <
phydro/L over the entire range of impact velocity. This reflects the strength of the
ceramics as discussed in [12].
The measured pprimary/L data can be well fit by a cubic equation in v:
B4C: pprimary/L = -1.213 + 2.178 v - 0.512 v2 + 0.044 v
3 (km/s) (7)
SiC: pprimary/L = 0.747 – 0.049 v + 0.185 v2 - 0.024 v
3 (km/s) (8)
AlN: pprimary/L = -1.258 + 1.842 v - 0.342 v2 + 0.022 v
3 (km/s) (9)
These cubic equations fit the data well [6-8] and are convenient but, of course,
they are purely empirical, do not asymptote to the hydrodynamic limit as v ,
and should not be used outside the range of the data.
Ceramic Armor Materials by Design 157
0.0
1.0
2.0
3.0
1.0 2.0 3.0 4.0 5.0
Impact Velocity, km/s
Pp
rim
ary
/L
Fig. 7. Primary penetration versus impact velocity for B4C
0.0
0.5
1.0
1.5
2.0
2.5
3.0
1.0 2.0 3.0 4.0 5.0
Impact Velocity, km/s
Pp
rim
ary
/L
Fig. 8. Primary penetration versus impact velocity for SiC.
0.0
0.5
1.0
1.5
2.0
2.5
3.0
1.0 2.0 3.0 4.0 5.0
Impact Velocity, km/s
Pp
rim
ary
/L
Fig. 9. Primary penetration versus impact velocity for AlN.
158 Ceramic Armor Materials by Design
Total Penetration, ptotal
Figures 10, 11, and 12 show the normalized total penetration, ptotal/L versus
impact velocity for the three ceramics. Again the horizontal dashed line is the
ideal hydrodynamic penetration. For each of the ceramics ptotal/L is less than the
hydrodynamic value up to impact velocities of about 4 km/s or higher.
Strictly speaking hydrodynamic penetration only applies to what is called here
primary penetration. As noted above pprimary/L is less than the hydrodynamic
penetration even at impact velocities as high as 4.6 km/s. A comparison of Fig. 7-
9 with 10-12 shows that except for a few of the lowest velocity tests ptotal/L >
pprimary/L. Total penetration is the sum of the primary penetration plus what has
been called “secondary penetration”, “residual penetration”, “after-flow”, or as
preferred here “Phase 3 penetration”, after Eichelberger and Gehring [13]. As
discussed by Orphal [14] Phase 3 penetration potentially involves two distinct
phenomena. The first phenomena is often called “after-flow” after Pack and
Evans [15] and later Tate [16] and is the extension of the target hole due to
momentum in the target material at the time the rod is fully eroded. The second
phenomena was called “secondary penetration” by Christman and Gehring [17]
and Allen and Rogers [18] and is the further penetration of the target by the
eroded rod debris in the case when p > t. Phase 3 penetration for these three
ceramics is discussed in some detail by Orphal [14 ] .
The measured ptotal/L data were also fit by cubic equations in v:
B4C: ptotal/L = -2.338 + 3.256 v - 0.821 v2 + 0.077 v
3 (km/s) (10)
SiC: ptotal/L = -1.438 + 1.904 v - 0.367 v2 + 0.030 v
3 (km/s) (11)
AlN: ptotal/L = -1.393 + 1.954 v - 0.365 v2 + 0.029 v
3 (km/s) (12)
The admonition above about the application of these purely empirical fits applies
to ptotal/L as well.
0.0
1.0
2.0
3.0
4.0
1.0 2.0 3.0 4.0 5.0
Impact Velocity, km/s
Pto
tal/L
Fig. 10. Total Penetration versus impact velocity for B4C.
Ceramic Armor Materials by Design 159
0.0
0.5
1.0
1.5
2.0
2.5
3.0
1.0 2.0 3.0 4.0 5.0
Impact Velocity, km/s
Pto
tal/L
Fig. 11. Total Penetration versus impact velocity for SiC.
0.0
0.5
1.0
1.5
2.0
2.5
3.0
1.0 2.0 3.0 4.0 5.0
Impact Velocity, km/s
Pto
tal/L
Fig. 12. Total Penetration versus impact velocity for AlN.
SIZE SCALING
These reverse ballistic experiments are small scale. The issue of scaling is best
addressed by performing full-scale tests, if possible, or at least larger scale tests.
Ten larger scale direct ballistic tests were performed against B4C targets. These
tests are described in detail in [6, 19]. The target design for these larger scale tests
was based on the reverse ballistic targets. The ceramic diameter was 87.5 mm. In
nine of these tests L/D = 13 tungsten alloy rods were impacted against the
confined B4C targets at 1.38 v 3.76 km/s. These rods had D = 4.22 mm in five
tests (1.38 v 2.96 km/s) and 2.92 mm in four tests (3.21 v 3.76 km/s). In
the tenth test (2.79 km/s) an L/D = 20, D = 2.79 mm rod was used. Thus in these
tests the ratio of ceramic diameter to rod diameter was 20-30. These tests were
essentially 5.5 times larger in scale than the reverse ballistic tests. Figure 13
compares the penetration velocity, u, measured in both the reverse ballistic tests
(Fig. 4) and these larger scale direct ballistic tests (labeled GRC). Figure 14
160 Ceramic Armor Materials by Design
compares ptotal/L for the two sets of experiments. The agreement between the data
from the small-scale reverse ballistic tests and the larger scale direct ballistic tests
is considered very good.
Seven similar larger scale direct ballistic tests were performed for AlN targets
by Piekutowski and Forrestal [20]. While not included here because of space
limitations the agreement between the reverse and larger-scale direct ballistic
experimental data for AlN is similarly very good [8].
0.0
1.0
2.0
3.0
4.0
1.0 2.0 3.0 4.0 5.0
Impact Velocity, km/s
Pen
etr
ati
on
Velo
cit
y, u
,
km
/s
GRC
Fig. 13. Comparison of penetration velocity with larger scale direct ballistic tests.
0.0
1.0
2.0
3.0
4.0
1.0 2.0 3.0 4.0 5.0
Impact Velocity, km/s
Pto
tal/L
GRC
Fig. 14. Comparison of total penetration with larger scale direct ballistic tests.
NORMALIZATION OF TOTAL PENETRATION FOR B4C, SiC AND AlN
For hydrodynamic penetration P/[L( p/ t)1/2
] = 1. To compare P/L for different
ceramics it is reasonable to attempt to normalize by the factor ( p/ t)1/2
. Figure 15
shows Ptotal/[L( p/ t)1/2
] versus velocity for B4C, SiC and AlN. As can be seen, for
these three ceramics this normalization approximately collapses the data to a
Ceramic Armor Materials by Design 161
single curve. Although not shown here, similar results are obtained for
Pprimary/[L( p/ t)1/2
].
Fig. 15. Ptotal / [L( p/ t)1/2
] versus impact velocity for B4C, SiC and AlN
0
0.2
0.4
0.6
0.8
1
1.2
1 2 3 4 5
Impact Velocity, km/s
Pto
tal/
[L*(
rho
p/r
ho
t)**
0.5
]
B4C
SiC
AlN
SUMMARY
Penetration depth and remaining rod length as functions of time were measured
in reverse ballistic tests for long tungsten rods impacting confined B4C, SiC and
AlN at velocities from 1.5 to 5 km/s. Penetration is steady state (or very nearly
steady state at the lowest impact velocities), i.e. penetration velocity = constant.
Penetration velocity and rate of rod erosion are both well described by linear
functions of impact velocity. Primary penetration is less than the ideal
hydrodynamic value over the entire range of impact velocity. Total penetration is
less than ideal hydrodynamic for impact velocities less than about 4 km/s.
Dividing penetration depth by the factor L( p/ t)1/2
nearly collapses the
penetration versus impact velocity data for the three ceramics to a single curve.
REFERENCES1V. Hohler and A. J. Stilp, “Hypervelocity impact of rod projectiles with L/D
from 1 to 32,” International Journal of Impact Engineering, 5, 323-331 (1987).2V. Hohler and A. J. Stilp, “Long rod penetration mechanics,” Chapter 5 in
High Velocity Impact Dynamics, Edited by Jonas A. Zukas. John Wiley, 1990.3B. R. Sorensen, K. D. Kimsey, G. F. Silsby, D. R. Scheffler, T. M. Sherrick
and W. D. deRosset, “High velocity penetration of steel targets,” International
Journal of Impact Engineering, 11, 107-119 (1991).4P. Lundberg, R. Renstrom and B. Lundberg, “Impact of metallic projectiles on
ceramic targets: transition between interface defeat and penetration,”
International Journal of Impact Engineering, 24, 259-275 (2000).
162 Ceramic Armor Materials by Design
5L. Westerling, P. Lundberg and B. Lundberg, “Tungsten long rod penetration
into confined cylinders of boron carbide at and above ordnance velocities,”
International Journal of Impact Engineering, 25, 703-714 (2001). 6D. L. Orphal, R. R. Franzen, A. C. Charters, T. L. Menna, and A. J.
Piekutowski, “Penetration of confined boron carbide at targets long rods at impact
velocities from 1.5 to 5.0 km/s,” International Journal of Impact Engineering, 19,
15-29 (1997). 7
D. L. Orphal and R. R. Franzen, “Penetration of confined silicon carbide
targets by tungsten long rods at impact velocities from 1.5 to 4.6 km/s,”
International Journal of Impact Engineering, 19, 1-13(1997). 8
D. L. Orphal, R. R. Franzen, A. J. Piekutowski, and M. J. Forrestal,
“Penetration of confined aluminum nitride targets by tungsten long rods at 1.5-4.5
km/s,” International Journal of Impact Engineering, 18, 355-368 (1996). 9
D. L. Littlefield, C. E. Anderson, jr., Y. Partom, and S. J. Bless, “The
penetration of steel targets finite in radial extent”, International Journal of Impact
Engineering, 19, 49-62 (1997). 10
C. E. Anderson, Jr., J. D. Walker, and T. R. Sharron, “The influence of edge
effects on penetration”, Proceedings: 17th
International Symposium on Ballistics
(Midrand, South Africa), pages 3-33 to 3-40, March 23-27, 1998 11
G. Birkoff, D. P. MacDougall, E. M. Pugh, and Sir. G. Taylor, “Explosives
with lined cavities,” Journal of Applied Physics, 19, 563-582 (1948). 12
C. E. Anderson, Jr., D. L. Orphal, R. R. Franzen, and J. D. Walker, “On the
hydrodynamic approximation for long rod penetration’” International Journal of
Impact Engineering, 22, 23-43 (1999). 13
R. J. Eichelberger and J. W. Gehring, Journal of the American Rocket
Society, 32, 1583-1591 (1962). 14
D. L. Orphal, “Phase three penetration,” International Journal of Impact
Engineering, 20, 601-616 (1997). 15
D. C. Pack and W. M. Evans, Proceedings Physical Society of London, B64,
298-302 (1951) 16
A. Tate, “A theory for the deceleration of long rods after impact,” Journal
Mechanics and Physics of Solids, 15, 387-399 (1967). 17
D. R. Christman and J. W. Gehring, “Analysis of high-velocity projectile
penetration mechanic,” Journal of Applied Physics, 37, 1579-1587 (1966). 18
W. A. Aleen and J. W. Rogers, “Penetration of a rod into a semi-infinite
target,” Journal Franklin Institute, 272, 275-284 (1961). 19
T. L. Menna, A. C. Charters, and A. J. Piekutowski, “Penetration
performance of confined boron carbide by continuous and segmented rods,”
General Research Corporation Report SB-90-0105 (1990). 20
A. J. Piekutowski and M. J. Forrestal, “Penetration into aluminum nitride
targets with L/D = 10 tungsten rods at impact velocities of 1.7, 2.2 and 2.7 km/s,”
Ceramic Armor Materials by Design 163
Report SAND91-0088.RS123/90/00007, Sandia National Laboratories,
Albuquerque, NM.
164 Ceramic Armor Materials by Design
DEPTH OF PENETRATION TESTING
Dr Bryn James
Defence Science and Technology Laboratories
Chobham Lane
Chertsey, Surrey, KT16 0EE
United Kingdom
ABSTRACT
The Depth of Penetration (DOP) test has been widely used for many years for
ranking the protective value of materials, most notably ceramics. This is
essentially a simple and straightforward test with definitive results. In practice
however, a significant number of factors must be taken into account to achieve
reliable and comparable results. Often, published results cannot be utilised to
augment other data sets as insufficient detail is given to allow the necessary
correction or normalisation to be made.
The aim of this paper is to provide details of the DOP measurement system
devised at the Defence Science and Technology Laboratories, Chertsey, UK and
to present the methodology for correction and normalisation of the data.
Guidelines will be given for choice of configuration of the target assembly.
INTRODUCTION
There are two principal methods by which a material may be tested
ballistically. The ballistic limit configuration reported in the literature (1) consists
of a relatively thin layer, or composite, system which is defeated by the
penetrator. Performance of this configuration is measured by the residual length
and velocity of the projectile, or by penetration of a witness pack. This
configuration is often used to investigate the effectiveness of specific armour
systems. Later configurations published in the literature are variations on the
semi-infinite backstop method first suggested by Bless and Rosenberg et al. (2,
3). Performance in this configuration is measured by the residual Depth of
Penetration (DOP) of the projectile into a backblock of reference material for
which the penetration depth of the projectile for direct impact is known. The
backblock dimensions are large ('semi-infinite'), such that the penetration is not
influenced by the proximity of edges or interfaces.
Ceramic Armor Materials by Design 165
To the extent authorized under the laws of the United States of America, all copyright interests in this publication are the propertyof The American Ceramic Society. Any duplication, reproduction, or republication of this publication or any part thereof, withoutthe express written consent of The American Ceramic Society or fee paid to the Copyright Clearance Center, is prohibited.
In order to assess the intrinsic performance of bulk material, the latter
configuration is preferred. Target dimensions are determined such that the
majority of the steady state penetration phase is accommodated within the
material under test. A small amount of this phase and the deceleration phase are
contained within the semi-infinite back-block to give a depth of penetration
measurement. Massive containment is often used for ceramic tiles to mimic the
effects of a laterally infinite target. This configuration is designed to avoid the
effects of elastic reflections from the lateral extents of the target assembly
affecting the penetrator, and to maintain impact induced pressure.
A problematical situation exists concerning medium scale, long rod type
penetration testing for which the backblock material is often rolled homogeneous
armour steel (RHA). Unfortunately the specification for RHA is different in
virtually every country, leading to many different values for ballistic efficiency
being quoted for the same material. We have investigated the use of alternative,
more universally available and better specified materials, but unfortunately, RHA
still seems to be the ideal material for this application. This situation also occurs
for other backblock materials for which an internationally agreed standard has not
been resolved.
Other factors that affect the derived value for ballistic efficiency include the
variation of depth-of-penetration with projectile yaw at impact, with projectile
velocity and of course with the overall target configuration.
BALLISTIC TESTING
In order to mimic the performance of a semi-infinite array of ceramic tiles, a
massive containment system was developed at Dstl, Chertsey (Figure 1).
Lateral containment
Section of rig showing
lateral and axial containment
Ceramic
RHA backblock
Figure 1. Ceramic containment system
166 Ceramic Armor Materials by Design
It should be noted that the containment ring is included to minimise the reflection
of impact induced stress waves from the periphery of the ceramic tile and to
maintain impact pressure. Any precompression of the ceramic has a negligible
effect upon the intrinsic ceramic performance as the small amount of pressure that
can be applied is negligible compared to the 5-20 GPa required to increase
significantly the failure strength of the material.
The corner of the containment ring, in our system, is relieved to allow the
ceramic tile to fit freely and to allow space for the introduction of a 1.0mm thick,
fully annealed, brass shim between the steel containment and the ceramic. This
shim has a similar acoustic impedance (18.2 MRayls) to that of steel
(25.4 MRayls), alumina (21.4 MRayls), boron carbide (22 MRayls) and silicon
carbide (25.2 MRayls) but being soft, will conform to any small irregularities in
the mating surface between the steel and the ceramic, providing an excellent
acoustic interface (Figure 2).
Figure 2 Brass shimming of ceramic within containment
The containment rig is assembled with the surface machined RHA (or
aluminium alloy) backblock between the rear of the ceramic tile (or tiles) and the
backplate of the rig. All bolts are tightened, in sequence, to a given torque.
IMPACT YAW CORRECTION
The penetration of a projectile subject to yaw at impact will be less than that
for an exactly normal impact. In order to allow for the experimental variation in
projectile yaw, depth-of-penetration measurements are corrected for yaw
according to the analysis of Bjerke et al (4). Note that in the following, 'yaw'
refers to the total yaw of the projectile, i.e. the angle between the longitudinal axis
of the rod and the velocity vector of the centre of mass.
Ceramic Armor Materials by Design 167
Bjerke’s correction factors are based upon the analysis of a very large number
of low yaw normal impacts of long rods into semi-infinite RHA over a range of
velocities from 1.2 to 4.7 kms-1
and a wide range of sizes and aspect ratios.
Upon impact, a penetration channel of diameter H is formed. If the projectile
yaw is such that the tail of the projectile of length L gouges the side of the
penetration channel, energy and mass of the projectile will be expended enlarging
the channel thus diminishing the total depth-of-penetration. As the penetration
channel is larger than the projectile diameter D, this condition will not occur until
a critical yaw angle has been reached. The critical yaw value cr is given by the
following relationship:
]2L / D)-(H[sin= -1
cr(1)
It is obvious from the above that the critical yaw value depends upon the
penetration channel diameter. It is not necessary to determine this value for each
impact as Bjerke et al. have performed an empirical fit to their large database of
impact geometry’s and velocities V. The penetration channel diameter may be
approximated by:
(2) 2V0.1286+V0.3388+1.1524=D / H
Where V is in kms-1
.
Bjerke calculated relative yaw (i.e. / cr) for a large number of impact
experiments so that an empirical fit could be made and the effective penetration as
a function of yaw could be calculated, eliminating the influence of velocity and
relative penetrator dimensions. Effective penetration factor, PENeff , is then given
by:
(3) ) / 11.46(cos=PEN creff
The yaw corrected penetration DOPyaw-corr may then be calculated from the
measured penetration DOPmeas using the following relationship:
(4) PEN / DOP=DOP effmeascorr-yaw
The original Bjerke analysis was performed for the impact of tungsten alloy
rods into steel. However, high intensity X-radiography has shown that the
dynamic crater diameter in ceramic materials is very similar to that in steel. We
168 Ceramic Armor Materials by Design
have also shown that use of the yaw correction significantly reduces the standard
deviation in the data set. It is therefore considered that this correction is valid.
VELOCITY NORMALISATION
The depth-of-penetration is a function of impact velocity. This function is
generally not linear over a wide range, and so a reference DOP assessment must
be made into the backblock material over the velocity range at which the material
will be tested. In order to obtain uniformity we must normalise both reference
DOP and residual DOP to a nominated test velocity.
If it is deemed necessary (i.e. in the case of large non-linearity within the
velocity range used), the full form of the DOP vs. velocity curve must be
determined, and a polynomial should be fitted. In our case, the relationship has
been found to be highly linear over a relatively large velocity range and so a
straight line fit may be used. The velocity normalised depth-of-penetration
DOPvel-norm at the reference velocity Vref is calculated from the measured depth-of-
penetration at the measurement velocity Vmeas by:
(5) )V-Vm.(+DOP=DOP measrefcorr-yawnorm-vel
Where m is the empirically derived slope of the DOP vs. velocity curve for
the specific penetrator.
CALCULATION OF CORRECTED AND NORMALISED DOP
The following details indicate the method for calculating the yaw corrected
depth of penetration and normalisation to a given reference velocity:
i/ Generate curve of DOP into the backing material vs. Velocity for the
velocity range required, using yaw corrected DOP's. Calculate straight line
regression or more complex function if required.
ii/ For each experiment, calculate yaw-corrected DOP.
iii/ Normalise yaw corrected DOP to Vref.
iv/ Determine DOPref at Vref using curve generated in i/.
Calculation of yaw-corrected DOP
i/ Calculate H/D at impact velocity using equation 2, hence find H.
ii/ Calculate cr at impact velocity using equation 1.
iii/ Calculate PENeff using equation 3 and the measured total yaw at
impact, hence find DOPyaw-corr using equation 4.
Ceramic Armor Materials by Design 169
FACTORS AFFECTING RESULTS
Rolled Homogeneous Armour Steel. The single most important factor affecting
the measured DOP results rests in the mechanical properties of the backing block.
The properties of RHA differ widely from nation to nation and in general the
specification for RHA is very wide. Often, the only reason for DOP results to be
valid is that, for economical reasons, the manufacturer supplies material with
properties as close as possible to the lowest limit of the specification. We have
found that, at the very least, a hardness measurement should be taken for every
batch of DOP material used, all backblocks must come from the same batch and
they must all be prepared in the same fashion.
During the course of experimentation at Dstl, Chertsey, (approximately 1000
DOP experiments), some interesting anomalies have been encountered in the
properties of RHA. Through thickness hardness transects generally show uniform
hardness in the bulk of the material with a slightly harder layer at the surface.
Occasionally, RHA plates have been seen with soft surface layers or with a
hardness gradient through the material. These plates appear visually to be the
same as standard plates, but may show a DOP up to 20% greater than standard
material. The importance of acceptance testing for each plate is apparent.
Ceramic Tile Size. Significant differences in the impact performance of a
given type of ceramic will be seen dependent upon the size of the sintered or
pressed tile. The furnace conditions necessarily must be different for different
sizes of tile, resulting in different residual stress states for different sized tiles. In
general, the larger the tile, the worse the ballistic performance, due to residual
stress in the material. We have encountered very large ceramic tiles with so much
residual stress that inadvertent damage during handling could result in
catastrophic fracture. To implement scaling experiments the ceramic material
should be cut from one large piece or large tiles should be machined to produce
smaller tiles. However, the surface state should be the same in each case.
Surface Preparation. The firing process for ceramic tiles often has an
advantageous effect on the surface, leaving a relatively flaw free layer with a
residual compressive stress. The presence of this layer tends to increase the
ballistic performance of the tile (up to ~5%). Often, this layer is machined away
in order to create a more precisely defined surface for interfacing with another
layer. It should be noted that any increases in performance gained by better
geometrical tolerancing may be offset by the removal of the beneficial surface
layer to be replaced by a surface microscopically damaged by the machining
process.
Lateral dimensions. In order for the lateral dimensions to have no influence
upon the measured DOP result, ideally the ceramic tile smallest lateral dimension
should be greater than 30 times the projectile diameter for relatively low velocity
impact. As the impact velocity increases, the lateral dimensions may become
170 Ceramic Armor Materials by Design
smaller. Above an impact velocity of 1600 ms-1
the tile size can be reduced until
at an impact velocity of 1800 ms-1
the lateral dimension of the tile may be only 15
times the projectile diameter. This should be considered a minimum. If no tiles of
adequate size are available, the tile may be clad in a supportive frame to mimic a
larger tile. Should a frame be fitted, the ideal minimum width of this frame should
be 0.5 x (30 x projectile diameter – the ceramic width). Of course, if the
calculated frame thickness is very thin, it may be omitted.
Ideally, any frame should be securely clamped to the ceramic with an acoustic
impedance matched soft shim between the ceramic and frame. It should be noted
that the lateral dimension guidelines are applicable in the case of a centre strike on
the target assembly. If there is appreciable shot dispersion from the launch
system, the target array should be larger to accommodate this dispersion.
RECORDING DATA
In order for published or recorded data to be valuable to other workers, the
following data should be reported:
Backing material Minimum: Manufacturer, type, hardness, density, axial and
lateral dimensions.
Desirable: Material model data.
Ceramic material Minimum: Manufacturer, type, density, axial and lateral
dimensions.
Desirable: Porosity, Material model data.
Projectile material Minimum: Manufacturer, type, density, dimensions.
Desirable: Hardness, strength, material model data.
Target configuration Minimum: All lateral and thickness dimensions and details
of any interlayers. Method of support.
Desirable: Tolerances and finish of mating surfaces,
tightening torques.
Impact Minimum: Impact velocity, yaw at impact, strike position.
Desirable: Yawing rate at impact.
CONCLUSIONS
The DOP test is a useful ranking test for ceramic materials. If attention is paid
to the detail and reproducibility of the target configuration, and if simple
corrections and a normalisation are carried out on the results, a significant amount
Ceramic Armor Materials by Design 171
of the inherent scatter in the data can be removed. The ranking measured in one
configuration is applicable to other configurations, even though the absolute
values of protective capability may not be the same for all configurations.
If a certain minimum data set is published for each experimental series, the
results may be easily utilised by other workers to supplement their own data sets.
ACKNOWLEDGEMENT
The work upon which this analysis is based was funded by the UK
Government Corporate Research Programme.
REFERENCES
1 J. A. Zukas, T. Nicholas, H. F. Swift, L. B. Greszczuk and D. R. Curran,
Impact Dynamics, J. Wiley and Sons, New York 1982
2 Z.Rosenberg, S.Bless, Y.Yeshurun, K.Okajima, “ A new definition of
ballistic efficiency of brittle materials based on the use of thick backing
plates”, in Impact loading and dynamic behaviour of materials Vol. 1 (ed.
C. Y. Chiem, H. D. Kunze and L. W. Meyer) pp. 491-498. DGM
Informationsgesellscaht mbH, Oberursel 1988
3 S.Bless, Z.Rosenberg, B.Yoon, “Hypervelocity penetration of ceramics”
Int.J.Impact Engng., 5 pp.165-171, 1987
4 T. Bjerke, G. Silsby, D. Scheffler, R. Mudd, “High yaw penetration
performance of long rod penetrators”, pp 191-198, Vol 3 13th
Int. Symp.
Ballistics, June 1992
172 Ceramic Armor Materials by Design
TRANSITION BETWEEN INTERFACE DEFEAT AND PENETRATION FOR
A GIVEN COMBINATION OF PROJECTILE-AND CERAMIC MATERIAL
Patrik Lundberg, René Renström and Lars Westerling
Swedish Defence Research Agency, FOI
Weapons and Protection
SE-147 25 Tumba, Sweden
ABSTRACT
At a certain impact velocity, the surface load generated by a projectile be-
comes critical and transition from interface defeat to penetration occurs. This
transition impact velocity is estimated by combining two models, one for the con-
tact load during interface defeat and one for the yield condition in the ceramic
target. The effects of the model parameters are studied with the aid of numerical
simulations, and the transition impact velocity is determined as a function of the
ceramic yield strength.
INTRODUCTION
By using devises for chock attenuation and load distribution in combination
with supporting confinement, it is possible to design ceramic targets capable of
defeating high velocity projectiles on the ceramic surface, so-called interface de-
feat or dwell [1-5]. The possibility to maintain interface defeat against long rod
projectiles for long interaction times (hundreds of s), which results in a nearly
static loading of the ceramic target material, has been shown experimentally [5].
The transition impact velocity, i.e., the impact velocity at which interface de-
feat ceases and penetration starts, has been determined experimentally for differ-
ent ceramic materials [4]. Also, analytical models for estimation of this velocity
have been derived [4,6]. The results of the experiments show that the transition
from interface defeat to penetration is distinct and related to the surface load.
In order to investigate the state of stress in the target material under conditions
of interface defeat, it is necessary to determine the contact pressure (negative of
normal stress at the contact interface) generated by the flowing projectile material.
Ceramic Armor Materials by Design 173
To the extent authorized under the laws of the United States of America, all copyright interests in this publication are the propertyof The American Ceramic Society. Any duplication, reproduction, or republication of this publication or any part thereof, withoutthe express written consent of The American Ceramic Society or fee paid to the Copyright Clearance Center, is prohibited.
This contact pressure and the corresponding state of stress is studied here by
means of continuum-dynamic simulations. First, a relation is derived for the
maximum contact pressure during interface defeat as a function of impact velocity
and projectile density, bulk modulus and yield strength. Then, using a simple con-
stitutive model for the ceramic material and the contact pressure distributions
from the simulations, a relation is established between the maximum contact pres-
sure and the state of stress in the target corresponding to incipient plastic yield
and large-scale plastic yield. With the aid of these relations, the transition impact
velocity is estimated as a function of the ceramic yield strength.
ANALYTICAL MODEL
In [4], analytical models were presented which make it possible to estimate
the maximum contact pressure generated by a long-rod projectile during inter-
face defeat and the corresponding state of stress in the ceramic target.0p
The projectile material was treated as linear elastic perfectly plastic, obeying
von Mises yield criterion. It was characterised by its density bulk modulus
and yield strength . The target surface was considered to be flat, rigid and
friction-free, and the axis of symmetry of the projectile was oriented perpendicu-
larly to the target surface.
,p
pK yp
It was shown in [4] that the maximum contact pressure at r = 0, z = 0 can
be expressed as0p
, (1) 0 1pp q
where2 2p p pq v , (2)
and is the impact velocity. The functionspv and , which represent
elastic and yield strength effects, respectively, were determined to be 1 2
and 3.27 , where
,p p yp pK q q . (3)
The radial distribution of the projectile load was approximated by one deter-
mined experimentally for a low-velocity water jet [7]. This load distribution, in
combination with Boussinesq’s elastic stress field solution [8] and a plastic slip-
line solution [9] for the indentation of a flat rigid die, led to the transition interval
174 Ceramic Armor Materials by Design
(4) 0low yc high ycp
for the transition contact pressure, where is the yield strength of the ceramic
in uniaxial compression. The coefficients and were determined in [4] to
be 1.47 (for Poisson’s ratio ) and 2.85, respectively. The ends of the in-
terval (4) correspond to incipient and large-scale plastic yield, respectively, of the
ceramic material. Because of relations (1) and (2) there is a corresponding interval
yc
low high
0.16
(5)low p highv v v
for the transition impact velocity. In Figure 1, the domain where yield occurs in
the ceramic target is illustrated correspondingly.
p lowv vincipient
plastic yield
low p highv v v
(a) (b)
p highv vlarge-scale
plastic yield
(c)
Fig. 1 Impact velocity (a) equal to the lower transition velocity , (b) in-
between the lower and higher transition velocity and (c)
equal to the upper transition velocity . The shape and location of the
domain where yield occurs is schematic.
lowv
highvlow pv v
highv
NUMERICAL SIMULATIONS
The AUTODYN code [10] was used for determining the distribution of con-
tact pressure and the resulting state of stress in a ceramic target during interface
Ceramic Armor Materials by Design 175
defeat. The simulations were two-dimensional with cylindrical symmetry. Two
types of simulations were performed. First, the contact pressure on a flat, rigid
and friction-free target surface was calculated using Eulerian simulations. A line-
arly elastic perfectly plastic constitutive model and von Mises yield criterion with
associated flow rule was used for the projectile material. A frictionless boundary
condition was used for the impact surface, and an inflow boundary condition was
used to simulate an infinitely long projectile. The simulations were performed
until a stationary contact pressure was reached. Different combinations of impact
velocities vp, bulk moduli Kp and yield strengths yp were used together with a
projectile density 17600 kg/mp3 in order to separate the influence of the com-
pressibility and the yield strength. By this technique, the functions and
were evaluated.
One of the contact pressure distributions was used in a second set of Lagran-
gian simulations in which the coefficients and were determined. It was
used as a boundary condition and the contact pressure was increased linearly in
order to follow the formation and growth of a plastic region, which finally
reached the contact surface. The target material was modelled as a linear elastic
perfectly plastic material with density kg/m
low high
3215c3, bulk modulus
GPa, Poisson's ratio and yield strength . A con-
tact pressure distribution corresponding to a Hertz indent [11] was also used for
comparison.
221cK 0.16 10 GPayc
RESULTS AND DISCUSSION
The maximum contact pressure p0 is shown in Table I for different combina-
tions of projectile bulk modulus Kp, yield strength yp and impact velocity vp.
Table I. Maximum contact pressure p0 for different combination of bulk
modulus Kp, yield stress yp and impact velocity vp.
p0 (GPa)Kp
(GPa)yp
(GPa) vp =1000 m/s vp = 1500 m/s vp = 2000 m/s vp = 2500 m/s vp = 3000 m/s
285 1 11.59 23.70 40.80 63.85 93.99
285 0.001 8.99 20.57 37.33 60.43 90.78
28500 1 11.63 23.04 38.67 58.49 82.87
28500 0.001 8.77 19.74 35.14 54.93 79.21
176 Ceramic Armor Materials by Design
The distribution of the contact pressure p and the normalised contact pressure
p/p0 for different impact velocities vp is shown in Figure 2. The curve for 1500
m/s in Figure 2(a) is the one used for the determination of the coefficients
and .low
high
The normalised contact pressure p/p0 for two different combinations of projec-
tile bulk moduli and yield strengths are shown in Figure 2(b). If the influence of
compressibility and yield strength is suppressed, corresponding to high bulk
modulus and low yield strength ( ), the contact pressure
distribution corresponds well to the one measured for a low velocity water jet [7].
On the other hand, low bulk modulus and high yield strength will give a narrower
pressure distribution as shown in the figure ( ). As a compari-
son, the pressure distribution for a Hertz indent [11] is also shown.
3240, 0.000114
32.4, 0.114
0 0.5 1 1.5 2 2.5 3
r/a
0
20
40
60
80
100
p (G
Pa)
1000 m/s
1500 m/s
2500 m/s
2000 m/s
3000 m/sp = 17600 kg/m 3
Kp = 285 GPa
yp = 1 GPa
0 0.5 1 1.5 2 2.5 3
r/a
0
0.2
0.4
0.6
0.8
1p/p
0Low velocity water jet
= 3240, = 0.000114
= 32.4, = 0.114
Hertz
(a) (b)
Fig. 2 (a) Contact pressure p and (b) normalised contact pressure p/p0 versus
normalised radius r/a. Filled circles in (b) represent the contact pressure
distribution from a low-velocity water jet [7], and dashed curve corre-
sponds to a Hertz indent [11].
The functions and obtained from the simulations are plotted in
Figure 3.
Ceramic Armor Materials by Design 177
0 0.05 0.1 0.15 0.2 0.25 0.3 0.35
-1
0
0.05
0.1
0.15
0.2
0 0.025 0.05 0.075 0.1 0.125 0.150
0.05
0.1
0.15
0.2
0.25
0.3
0.35
0.4
1 1.93 23.62 7.04
(a) (b)
Fig. 3 The contribution of (a) compressibility versus for small values of
(< 0.000114) and (b) material strength versus for large values of
(> 360). Filled circles are simulations and the solid curves are least square
fits to the data.
1
The coefficients and for incipient and large-scale plastic yield are
shown in Table II.low high
Table II. The coefficients and .low high
Hertz indent Projectile load
Simulation Data in [4] Simulation
low 1.46 1.47 1.48
high - 2.85 2.73
From least square fits to the data in Figure 3 together with the data from Table
II, the estimates
1 1.93 , 23.62 7.04 , , (6) 1.48low 2.73high
based on the simulations are obtained.
Relations (1) to (6) give an interval for the transition impact velocity v as a
function of the ceramic yield strength. This interval is shown in Figure 4 together
with experimental data for silicon carbide [4]. The transition impact velocity for
silicon carbide determined experimentally is between 1645 m/s and 1705 m/s [4]
and the maximum yield strength, determined from plate impact experiments, is in
p
178 Ceramic Armor Materials by Design
the range of 12.5-14.5 GPa [12,13]. The density, bulk modulus and yield strength
used for the projectile material was =17600 kg/mp3, 285 GPa and
GPa, respectively.pK
1.2yp
10
yc
yc
lowv
1.49low
high
0 51000
1250
1500
1750
2000
v p
(m/s
)
SiC
Tungsten projectile
vlow
vhigh
15 20 25
(GPa)
Fig. 4 Transition velocity versus ceramic yield strength . The curves for
and correspond to incipient and large-scale plastic yield.pv
highv
The values in Table II for the coefficients and differ from the ones
obtained in [4].low high
The contact pressure distribution for a low-velocity water jet [7] used in [4],
gave a slightly lower value of than the projectile pressure distribution used
here, the curve for 1500 m/s in Figure 2(a). The factor can be solved analyti-
cally with Boussinesq’s elastic stress field solution [8] for arbitrary load distribu-
tions. This method gives for the Hertz pressure distribution and
for the projectile pressure distribution.
low
w
low
1.46lo
The coefficient corresponds to large-scale yield, viz., the instant when
the plastic region beneath the projectile reaches the surface and penetration takes
place. In the simulations, this coefficient was determined when the plastic region
reached the target surface on the axis of symmetry. This evaluation method gave a
slightly lower value of
high
high than that obtained in [4]. The differences can be re-
lated to the use of different material models and contact loads. In [4], the coeffi-
cient was based on the solution of a flat rigid cylindrical die indenting a
semi-infinite rigid-plastic medium [9], while an elastic perfectly plastic constitu-
tive model was used in the simulations together with a bell-shaped surface load.
Values of this coefficient between 2.5 and 3.0 have been reported from simula-
tions of the total loading history when a projectile starts to penetrate a target [14].
Ceramic Armor Materials by Design 179
CONCLUSIONS
A relation between the yield strength of the ceramic material and the transi-
tion impact velocity with regard to interface defeat has been obtained from nu-
merical simulations. This relation provides two limits, one lower, corresponding
to incipient yield in the target material, and one higher, corresponding to large-
scale yielding. These limit values agree well with analytical results published ear-
lier [4]. The transition impact velocity obtained experimentally for silicon carbide
falls in-between these two limits.
REFERENCES
[1] G. E. Hauver, P. H. Netherwood, R. F. Benck and L. J. Kecskes. Ballistic
performance of ceramic targets. Army Symposium On Solid Mechanics.
USA 1993.
[2] G. E. Hauver, P. H. Netherwood, R. F. Benck and L. J. Kecskes. Enhanced
ballistic performance of ceramic targets. 19th
Army Science Conference.
USA 1994.
[3] E. J. Rapacki, G. E. Hauver, P. H. Netherwood and R. F. Benck, Ceramics
for armours- a material system perspective. 7th
Annual TARDEC Ground
Vehicle Survivability Symposium. USA 1996.
[4] P. Lundberg, R. Renström, B. Lundberg. Impact of metallic projectiles on
ceramic targets: transition between interface defeat and penetration. Int J
Impact Engng 2000;24:259-275.
[5] P. Lundberg, R. Renström, L. Holmberg. An experimental investigation of
interface defeat at extended interaction times. Proc 19th Int Symp on Ballis-
tics, Switzerland: 2001;3:1463-1469.
[6] LaSalvia JC, Horwath EJ, Rapacki EJ, Shih CJ, Meyers MA. Microstruc-
tural and micromechanical aspects of ceramic/long-rod projectile interac-
tions: dwell/penetration transitions. Fundamental Issues and Applications of
Shock-Wave and High-Strain-Rate Phenomena, Staudhammer KP, Murr
LE, Meyers MA, Elsevier Science, pp 437-446, 2001.
[7] Reich F. Omlenkung eines freien Flussigkeitsstrahles an einer zur
Strömungsrichtung senkrecht stehenden ebenen Platte. Diss Hannover:
1926, (oder VDI-Forsch. –Heft 290).
[8] Y. C. Fung. Foundations of solid mechanics. Prentice-Hall, 1965.
[9] R. T. Shield. On the plastic flow of metals under conditions of axial symme-
try. Proc R Soc, A, 233, 267, 1955.
[10] K. Birnbaum, M. S. Cowler, M. Itoh, M. Katayama and H. Obata,
AUTODYN - an interactive non-linear dynamic analysis program for micro-
computers through supercomputers. Ninth International Conference on Struc-
tural Mechanics in Reactor Technology. Lausanne, Switzerland, (1987).
180 Ceramic Armor Materials by Design
[11] K. L. Johnson, Contact Mechanics, Cambridge University Press, 1985.
[12] N. Bourne, J. Millett, I. Pickup. Delayed failure in shocked silicon carbide.
J Appl Phys. 81(9), 1 May 1997.
[13] R. Feng, G. F. Raiser and Y. M. Gupta. material strength and inelastic de-
formation of silicon carbide under shock wave compression. J Appl Phys.
83(1), 1 January 1998.
[14] Z. Rosenberg, E. Dekel. Material similarities in long-rod penetration me-
chanics. Int J Impact Engng 2001;25:361-372.
Ceramic Armor Materials by Design 181
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Shock and High Strain Rate Dynamic
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DYNAMIC FRACTURE OF CERAMICS AND CMC
Albert S. Kobayashi
University of Washington
Department of Mechanical Engineering
Seattle, Washington 98195-2600
ABSTRACT
This paper reviews the limited literature on dynamic fracture mechanicscharacterization of ceramics and ceramic matrix composites (CMC). Dynamic
fracture toughness, KId, at room and elevated temperature of reaction bonded and
hot-pressed Si3N4, Al2O3, partially stabilized zirconia (PSZ), TiB2-particulate
reinforced SiC (TiB2p/SiC), and SiC-whisker reinforced Al2O3 (SiCw/Al2O)) are
presented. Dynamic stress intensity factor, KID versus crack velocity relations at
room and elevated temperature of Al2O3, SiCw/Al2O3, PSZ and Si3N4 are also
discussed. Dynamic crack arrest stress intensity factor, KIa, was only detected in
PSZ.
INTRODUCTION
Dynamic fracture mechanics encompasses the three phenomena of dynamic
crack initiation, i.e. crack initiation under dynamic loading, rapid crack
propagation, and arrest of a rapidly propagating crack. While early papers on
dynamic fracture mechanics dates back to the 1950’s1-4
, studies on dynamic
fracture mechanics started in the 1970's with the need to predict the extent of rapid
crack propagation in a nuclear power pressure vessel under emergency core
cooling and the effectiveness of a crack arrester in a large marine structure. As a
result of such concerted efforts, much is known on the dynamic response of a
rapidly propagating crack in metals and polymers. Unfortunately, the same cannot
be said about ceramics and CMC due to their extremely low static initiation
fracture toughness, i.e. KIC. With no tough ceramics in sight, design of a safe-fail
ceramic components is based on avoiding fracture all together or to promote the
use of ceramic components as a one-time energy absorber through fragmentation.
Both applications circumvent research in dynamic fracture mechanics of ceramics
and CMC.
Ceramic Armor Materials by Design 185
To the extent authorized under the laws of the United States of America, all copyright interests in this publication are the propertyof The American Ceramic Society. Any duplication, reproduction, or republication of this publication or any part thereof, withoutthe express written consent of The American Ceramic Society or fee paid to the Copyright Clearance Center, is prohibited.
In the following sections, a cursory review of the state of science of dynamic
fracture mechanics will be given. Procedures for dynamic fracture mechanics
characterization and properties peculiar to ceramics and CMC will be discussed.
HISTORICAL REVIEW
The early papers in dynamic fracture mechanics were simple extensions of
Griffith's instability criterion for predicting the onset of crack propagation. Mott1,
Roberts and Well3
and Berry4 added varying forms of an estimated kinetic energy
rate term to Griffith's balance of energy rate equation to account for the global
kinetics associated with a moving crack. This approach did not account for the
dynamic crack-tip state of stress and the possible difference in the static and the
dynamic fracture processes.
The moving Griffith's crack, which was derived by Yoffe2
during this early
period, did provide a crack velocity independent stress intensity factor for a crack
velocity dependent crack tip stress field. Using her solution, Yoffe predicted a
crack kinking angle of about 63o at a crack velocity of about sixty percent of the
shear wave velocity thus leading to a crack branching criterion which depended on
a critical crack velocity. While Yoffe's solution was a historical first, the anomaly
of her modeling resulted in an infinite energy release rate as the crack velocity
approached the Rayleigh wave velocity. Subsequent solutions for a constant
velocity crack initiating from zero and finite crack lengths by Broberg5 and
Baker6, respectively showed that the energy release rate approached zero as the
crack velocity approached the Raleigh wave velocity. The corresponding crack tip
stress fields were also characterized by crack velocity dependent stress intensity
factors.
Early views on crack arrest considered the arrest to be an inverse of the onset
of crack propagation7, namely that a propagating crack would arrest when the
instantaneous static stress intensity factor KI < KIC. Many tests and research
programs were conducted to verify or discredit this postulate with raging
controversies at times on the physical significance of dynamic crack arrest stress
intensity factor, KIa. Ensuing experimental8 and numeical
9 analyses, however,
suggested that KIa is a material property and that static analysis is not sufficient for
predicting the arrest of a propagating crack.
FUNDAMENTAL EQUATIONS IN DYNAMIC FRACTURE
Ceramics exhibits cleavage fracture at room as well as at elevated
temperature. This is fortunate since most of the theoretical developments in
dynamic fracture are confined to linear elastic fracture mechanics (LEFM).
However, the additional complexities of fiber pullouts and fracture involved in
dynamic fracture of CMC are yet to be addressed. Available theoretical solutions
in dynamic fracture are few, some of which are discussed in the following
186 Ceramic Armor Materials by Design
sections, and are limited to a self-similar crack extending at a constant velocity in
an infinite solid. Despite these limitations these solutions can be used to deduce
the crack tip state of stress as well as to extract the dynamic stress intensity factor.
Stationary Crack Impacted by a Tension Wave
The dynamic initiation stress intensity factor, KId, of a stationary semi-infinite
crack, which is impacted by square plane tension wave of duration t, in an infinite
solid was given by Freund10
. For a ramp tensile pulse loading, KId is a simple
superposition of the discrete KId values of the corresponding incremental plane
tension waves. Unlike its static counterpart, this KId does not involve a
characteristic length dimension. If, however, the crack starts to propagate rapidly
after an incubation time, then the resultant stress intensity factor, KID, is modified
by a scalar function of the crack velocity.
Crack Propagating at Constant Velocity
The state of stress at the tip of a crack propagating at a constant velocity in a
two dimensional, isotropic, homogeneous elastic material has been derived by
Nishioka and Atluri11
who provided the asymptotic crack tip stress and
displacement fields in infinite series. The singular, first order term in the infinite
series with a dynamic stress intensity factor, KID, is the most significant term in
the crack tip stress field. In addition, the second order term has been shown to
govern crack kinking and branching angle12,13
.
DYNAMIC INITIATION FRACTURE TOUGHNESS
Dynamic initiation fracture toughness, KId, is commonly determined by impact
loading a ceramic or CMC fracture specimen by a split Hopkins bar or by a drop-
weight.
Split Hopkinson Bar Tester
The split Hopkinson bar tester, which has been used extensively for impact
testing of metals and ceramics, was modified by Duffy et. al.14,15
, as shown in
Figure 1, to impart tension directly to the specimen without a prior compression
wave. The compressive wave developed by an explosive charge is shaped into a
tensile pulse through multiple reflections and then propagates down the steel bar.
Duffy et al14,15
determined the dynamic initiation fracture toughness, KId, of
precracked Al2O3 and SiCw/Al2O3 bar specimens. Table I shows their Al2O3
results at room and elevated temperature testing. By adding a pre-torque to the
bar specimen, the test setup was also used to measure the fracture toughness under
combined modes I and III fracture 16
.
Ceramic Armor Materials by Design 187
Figure 1 Split Hopkinson bar test for dynamic fracture testing14,15
.
Table I. Dynamic and static initiation fracture toughness of Al2O314,15
.
Temp KId KIC KId/KID
(oC) (MPa m
1/2) (MPa m
1/2)
20 3.5 2.7 1.3
900 3.4 2.2 1.5
1100 3.1 2.2 1.4
1300 2.0 1.4 1.4
Drop Weight Test
An early study on KId of ceramics was based on a static evaluation of the
impact data obtained from drop-weight loaded, single edge-notched (SEN), three-
point bend (TPB) specimens17
. This static analysis was subsequently replaced by
a dynamic finite element (FE) analysis of a pre-cracked21
, SEN TPB specimen to
which the impact load and the crack extension histories were prescibed18-20
. As
the load was measured outside the furnace for elevated temperature testing, the FE
model also included the impact rod in its load train. The crack extension history
was monitor by a calibrated laser interferometric displacement gage system22
.
The KId of Al2O3, TiB2p/SiC and SiCw/Al2O3 CMC thus obtained are listed in
Table II.
Table II Dynamic and static initiation fracture toughness of Al2O320
and CMC19
.
Mat’l Temp KId KIC KId/KIC
(oC) (MPa m
1/2) (MPa m
1/2)
Al2O3 20 5.7 4.3 1.3
1000 4.3
TiB2p/SiC 20 5.7 5.2 1.1
1000 5.1
SiCw/Al2O3 20 6.2 6.2 1.0
1000 6.1
188 Ceramic Armor Materials by Design
Similar drop weight loading system and SEN-TPB pre-cracked21
specimens
were used to determine KId of PSZ and Si3N4 through a temperature of
1200oC
23,24. Caustics method combined with an ultra-high speed camera was
used to determine KId as well as KID during rapid crack propagation. These results
will be discussed together in a later section on dynamic crack propagation.
A novel variation of the drop weight testing is the one-point bend (OPB), pre-
cracked21
specimen which is suspended by ceramic threads in an infrared image
furnace25,26
. Since the thin threads are broken at the instant of impact, the
specimen breaks without any constraint at the two end supports. A pair of
semiconductor strain gages on the impact rod was used to measure the impact
force. KId was obtained from the equation of motion of the impacted OPB
specimen. The KId of five ceramics at room temperature and 1200oC are shown in
Table III. The KId rate was about 1.2 x 105 MPa m
1/2/sec. KId of SiC, Si3N4 and
Al2O3 remained constant but KId of PSZ decreased substantially at 600oC.
Table III Dynamic and static initiation fracture toughness of ceramics24
.
Ceramics Temp KId KIC KId/KID
(oC) (MPa m
1/2) (MPa m
1/2)
SiC 20 6.3 5.5 1.2
1200 5.9
Si3N4 20 6.0 6.0 1.0
1200 6.0
PSZ 20 7.0 4.0 1.8
600 3.0
Al2O3 20 5.2 4.5 1.2
1200 5.2
Al2O3/ZrO2 20 10.2 6.5 1.6
Instrumented Charpy Impact Test
Quasi-dynamic analysis of an instrumented Charpy impact test has been used
by T. Kobayashi et. al.27,28
for KId determination at room temperature. Unlike the
other results, their KId of Al2O3 and Si3N4 remained essential constant with
increasing KId rate and then suddenly increased at a KId rate of 105 MPa m
1/2/sec
as shown in Figure 2.
DYNAMIC FRACTURE OF CERAMICS AND CMC
Impact failures of ceramics and CMCs are characterized by shattering which is
a complex phenomenon involving a multitude of simultaneous micro-crack
generation, growth and coalescence into macro-cracks which in turn grow, branch
and coalesce. Intact fibers in CMC do not necessarily arrest a propagating crack
in the brittle ceramic matrix, as the propagating crack is known to tunnel around
Ceramic Armor Materials by Design 189
Figure 2 Effect of KId rate on KId28
.
the fibers with little crack opening. Theoretically, a CMC could have a larger KId
and a dynamic crack arrest stress intensity factor, KIa, in order to resist rapid crack
propagation. Once the laws governing a single crack is known, a statistical or a
fractal analysis of the many branched cracks and the laws governing fiber
fracture/pull-out can be used to predict the overall dynamic response of the
impacted structural component.
Dynamic fracture mechanics study of a rapidly propagating crack in ceramics
and CMC, however, are virtually non-existent except for the papers by Shimizu et.
al.23,24
and the author and his colleagues18-20
. The experimental procedures used
were discussed in the previous section and thus only the results are presented in
the following.
Figure 3 shows the resultant crack velocity versus the KID relation for Al2O3
where little differences are noted between the data of room temperature and
1000°C. If the cluster of data at the left end did not exist, then the well-known
gamma shape curve, which has been observed in metals and polymers, could have
been obtained. Figure 3, however, shows that the crack continues to propagate
slowly, i.e. at speeds ranging from 10 to 40 m/s under a KID less than the KIC and
is consistent with previous findings29
. Also shown in this figure is the KID versus
190 Ceramic Armor Materials by Design
crack velocity relation for statically loaded specimens under fixed displacement
loading at room temperature.
Figure 3 Crack velocity versus KID of Al2O319
.
Figure 4 shows the resultant crack velocity versus KID relations of TiB2p/SiC
CMC impacted at room temperature and 1200°C. The crack velocity under
impact loading is relatively constant during the entire crack propagation history.
Also shown is the crack velocity versus KID relation for a statically loaded
specimen.
Figure 4 Crack velocity versus KID of TiB2p/SiC19
.
Ceramic Armor Materials by Design 191
Figure 5 shows KID versus the resultant crack velocity of PSZ impacted at
room temperature. KIa, which was not observed by others, was obtained after the
crack had propagated about 100 m/sec and arrested at Kia = 4.0 MPa m1/2.
.
Figure 5 KID versus crack velocity relation of PSZ23
.
Figure 6 shows the KID versus crack velocity relation of Si3N4. The crack
velocity observed in Figures 5 and 6 are order of magnitude higher than those of
Figures 3 and 4, possibly due to the sharp but machined notch tip, which obliterate
the trailing fracture process zone of a real crack30
and yielded a larger KId and
hence higher stored energy to drive the crack. While the trend of a decreasing
crack velocity with decreasing KID is observed, the available data in Figures 2, 4
and 5 do not indicate the existence or lack of existence of a KIa in these ceramics
and CMC which lack the stress induced transformation of PSZ. Crack arrest,
however, has been observed in chevron-notched, three point bend specimens,
which were machined from the same SiCw/Al2O3 ceramic composites and which
were loaded under an extremely small displacement rate of 0.01 mm/min29
. The
run-arrest events in this test were characterized by small crack jumps of about 0.8
mm, which initiated at the sharp crack tip in the chevron notched specimens. Once
the excess driving force had been dissipated during rapid crack propagation under
static loading and the crack had entered a region of KID < KIC, crack arrest was to
be expected. Figures 3 and 4 show that such was not the case.
192 Ceramic Armor Materials by Design
Figure 6 KID versus crack velocity relation of Si3N424
.
FRACTURE MORPHOLOGY
The lack of crack arrest was then attributed to the difference in the fracture
morphologies of extremely slow and rapid crack extensions. This postulate was
tested by extensive fractography analysis of the statically and impact loaded Al2O3
specimens31
. While intergranular fracture was the dominant failure mode in both
specimens, some transgranular fracture was observed in all regions of the fracture
surface. The percentage areas of transgranular fracture decreased from an average
of 16% during the initiation phase to an average of 10% at slower crack
propagation in the impacted specimen. For the statically loaded specimen, the
percentage of transgranular areas decreased from 5 to 2%. The higher percentage
areas of transgranular fracture during the initiation phase can be attributed to the
higher crack velocity and the higher KID due to the overdriving force generated by
the blunt crack tip. Fractography analysis also showed that rapid crack
propagation is always accompanied with transgranular fracture regardless of the
magnitude of the driving force, i.e., KID and the crack velocity. In contrast, the
fracture morphology for stable crack growth showed the dominance of
intergranular failure. The percentage area in excess of 10% at the lowest KID of
about 1.5 MPa m1/2
was obtained from the impacted specimen. This data suggests
that the continuous input of work during the fracture process generated a higher
percentage area of transgranular failure with little chance of crack arrest.
The failure energy of a single crystal ceramic, i.e., energy required for
transgranular fracture, is generally higher than that of a polycrystalline ceramic
thus suggesting that transgranular failure requires more energy than intergranular
Ceramic Armor Materials by Design 193
failure. Transgranular failure thus provides a larger driving force but also a
competing higher resistance.
These results on Al2O3 showed that under rapid crack propagation,
transgranular fracture does occur both at high as well as at a low. KID was
associated with the corresponding high and low crack velocities, respectively.
However, the kinematic constraint of a rapidly extending flat crack front must
have enforced a locally moderate transgranular failure and drove the crack at a
lower KID, thus reducing the chance for crack arrest even at KID < KIC. A low
percentage area of transgranular failures, i.e. 2%, thus continued to drive the crack
at a subcritical KID.
CONCLUSIONS
The paucity of data is symptomatic of the developing state of this field where
few laboratories throughout the world have the resources to undertake dynamic
and impact characterization of ceramic composite at an operating temperature in
excess of 1200o C. Thus, much of the actual test data was limited to room
temperature testing of Al2O3 and Si3N4.
The limited data suggests that:
1. Dynamic initiation fracture toughness, KId, of most ceramics and CMC is
slightly larger than the static initiation fracture toughness, KIC.
2. Significant increase in KId of PSZ at room temperature was observed.
Likewise, the KId of PSZ at 600oC decreased significantly.
3. For a rapidly propagating crack in PSZ at room temperature, the dynamic
crack arrest stress intensity factor, KIa, was almost equal to its KIC.
ACKNOWLEDGMENT
The author gratefully acknowledges the financial support of the Office of
Naval Research through ONR Contract N00014-87-K-0326 through which many
of the results reported in this paper were generated.
194 Ceramic Armor Materials by Design
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15S. Suresh, T. Nakamura, Y. Yeshurun, K.-H. Yang and J. Duffy, "TensileFracture Toughness of Ceramic materials: Effects of Dynamic Loading andElevated Temperatures," Journal of American Ceramic Society, 73 (8), 2457-66(1990).
16S. Suresh and E. K. Tschegg, "Combined Mode I-Mode III Fracture ofFatigue Precracked Alumina," Journal of American Ceramic Society, 70 (10) 726-733 (1987).
17S.T. Gonczy and D.L. Johnson, "Impact Fracture of Ceramics at HighTemperature," Fracture Mechanics of Ceramics, 3, edited by. R.C. Bradt, D.P.H.Hasselman and F.F. Lange, Plenum Press, 495-506, New York (1978).
18A.S. Kobayashi, A.F. Emery, and B.M. Liaw "Dynamic Fracture Toughnessof Reaction Bonded Silicon Nitride," Journal of American Ceramic Society, 66 (2),151-155 (1983).
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20Y. Takagi and A.S. Kobayashi, ”Further Studies of Dynamic FractureResponses of Alumina and SiCw/Al2O3 Composite,” Proceedings of the
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27T. Kobayashi, K. Matsunuma, H. Ikawa and K. Motoyoshi, “Evaluation ofStatic and Dynamic Fracture Toughness in Ceramics,” Engineering FractureMechanics, 31 [5], 873-885 (1988).
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196 Ceramic Armor Materials by Design
COMPRESSIVE FRACTURE OF BRITTLE SOLIDS UNDER SHOCK-WAVE
LOADING
G. I. Kanel S. J. Bless
Institute for High Energy Densities The University of Texas at Austin
IVTAN, Izhorskaya 13/19, Institute for Advanced Technology
Moscow, 127412 Russia 3925 W. Braker Lane, Suite 400
Austin, Texas 78759
ABSTRACT
The behavior of different kinds of brittle materials, including single crystals,
glasses, and ceramics, under shock wave loading (uniaxial strain conditions) and
impact loading under uniaxial stress conditions, is reviewed and compared from
the viewpoints of mechanisms and criteria of plastic deformation and compressive
fracture.
GENERAL BEHAVIOR OF BRITTLE MATERIALS UNDER COMPRESSION
Mechanisms of inelastic deformation of brittle materials under compression
were initially investigated for rocks (see review papers1,2,3
). It was found that
fracture under one-dimensional stress conditions or at relatively low confining
pressure often occurs by axial splitting. For greater confining pressures, failure
occurs by shear faulting at an angle less than 45˚ to the loading axis. Extensive
compressive fracture is preceded by microcracking. Orientations of microcracks
are predominantly within 10˚ of the direction of compression.4 Crack density
increases as macroscopic deviatoric stress increases above a distinct threshold
level. Faults and other macroscopic fractures appear to form after attainment of
the ultimate compressive stress, which is called the failure stress. Beyond the
point of peak load, the failure becomes unstable. In the post-failure region of
compression, the load-carrying capacity drops rapidly to a low value.
Since cracks occupy volume, their formation is accompanied with a decrease
in the average matter density. This nonlinear inelastic volume change is
commonly referred to as dilatancy or bulking.5 The bulking effect grows with
increasing deviator stress and decreases under confining pressure. Typically the
onset of the dilatation region occurs between one-third and two-thirds of the
failure stress. Unloading from this stress region yields a permanent residual
volume increment.
Ceramic Armor Materials by Design 197
To the extent authorized under the laws of the United States of America, all copyright interests in this publication are the propertyof The American Ceramic Society. Any duplication, reproduction, or republication of this publication or any part thereof, withoutthe express written consent of The American Ceramic Society or fee paid to the Copyright Clearance Center, is prohibited.
The dilatancy is accompanied by a hysteresis in physical properties between
loading and unloading which is manifested mostly in the lateral strains, while the
axial strain is nearly elastic and almost completely recoverable. The lateral
dilatantional strains are attributed to opening of axial cracks. Formation of open
axial cracks at a fraction of the maximum stress is also suggested by variations of
sound velocity in axial and radial directions: velocity in the axial direction is
hardly changed by stress, whereas the sound velocity in the radial direction begins
to decrease at about half the failure stress and may drop 10 to 20%.5
A confining pressure strongly affects the strength and inelastic behavior of
brittle materials. The deviator stresses at which microcracking starts or failure
occurs increase as the confining pressure increases. At a sufficiently high pressure
a transition from brittle to ductile response usually occurs. For example, Heard
and Cline 6 observed failure followed immediately after essentially elastic
deformation of Al2O3, AlN, and BeO when the confining pressure was low, but
there was a transition to more ductile response at high pressures. The ultimate
compressive strength of ceramics increases rapidly with pressure below the brittle-
ductile transition; above this threshold, the ultimate strength is nearly constant.
These ceramics also exhibit increasing ductility when the confining pressure is
above the brittle-ductile transition. The pressure of transition from brittle to
ductile response is different for different materials. Alumina, for example,
remains brittle at confining pressures at least up to 1.25 GPa. However, extensive
evidence of ductility by both slip and twinning was observed in alumina and
sapphire at room-temperature indentation deformation.7,8
POSSIBLE MECHANISMS OF MICROCRACKING UNDER COMPRESSION
Open cracks, like other voids, may nucleate and grow only when at least one
principle stress is tensile. Even in the case of overall compressive loading, small
regions of tension may appear inside the body as a result of modification of the
stress field by different concentrators, such as grain boundary contacts,
microcracks, and cavities in the incident materials, etc.9,10
Hence, cracks may
grow in response to this local tensile stress.
Intuitively, it appears that if easy shear is allowed within a limited band with
fixed tips inside a body, rarefaction and compression regions will be created near
the tip, as illustrated schematically in Fig. 1. There should also be concentration of
shear stresses in the crack plane ahead of the tip. The rarefaction may initiate a
tensile crack which can grow in the direction perpendicular to maximum tension
out of the crack plane, whereas localized shear may propagate further in the crack
plane. For different materials and various loading rates, Kalthoff observed that
both of these modes of shear failure initiated at the crack tip.11
198 Ceramic Armor Materials by Design
Rarefaction
Compression
Shear band
Tensile
crack
Figure 1. Schematic of the failure initiation at mode-II crack tip.
Griffith postulated that isotropic materials contain randomly oriented flaws or
cracks in all directions which significantly alter the stress field within the
material.9
The basic hypothesis of Griffith’s model is that fracture occurs when
the most vulnerably oriented crack begins to extend under applied stress. The
extension of the crack is assumed to occur when the maximum tensile stress
component at any point around the crack reaches the critical value needed to
overcome the interatomic cohesion of the material. The Griffith theory, or at least
its basic premise that fracture starts from flaws, is fundamental to all
investigations of brittle fracture.
Brace and Bombolakis observed the growth of cracks in glass and polymer
plates under compression.12
They found that the most severely stressed cracks
were inclined at about 30˚ to the axis of compression. The cracks, when either
isolated or placed in an array, grow along a curved path that becomes parallel with
the direction of compression. When this direction is attained, growth stops. The
resultant kinked crack consisted of a central crack with sliding surfaces, which is
inclined to the direction of compression, plus two cracks emanated from its ends,
which are called “wing cracks”. Modern theories of brittle fracture and dilatancy
under compression are mostly based on the development of the wing crack model.
More recently, a series of similar experiments was performed by Nemat-
Nasser and Horii.13,14
They also have shown that the relative sliding of the faces
of one or even an array of pre-existing cracks leads to the formation of tension
cracks which grow in the direction of maximum compression. A lateral
compression reduces the final crack length, whereas even small lateral tension
increases it. In the model experiments with specimens containing a number of
Ceramic Armor Materials by Design 199
randomly oriented cracks or rows of inclined coplanar cracks, axial splitting,
rather than localized shear failure, was observed.
Thus, in both natural and man-made materials, crack growth is directed
preferably along the compression and does not immediately produce a mechanical
instability as it does in tension. The tensile axial cracks which may open and grow
at stress concentrators under overall compression cannot be considered
themselves as a mechanism of an inelastic shear strain, but they may facilitate
shear and rotation of blocks of matter relative to each other and in this way to
contribute to deformation. The stress required to cause additional crack extension
increases after some crack growth has been initiated. Macroscopic faults are
formed out of systems of cracks.
DYNAMIC STRENGTH PROPERTIES OF SINGLE CRYSTALS AND
GLASSES
First consider plate impact experiments on sapphire and ruby samples backed
by water, in which the rear surface velocity histories were measured.15,16
Results
are shown in Fig. 1. These data exhibit most of the peculiarities of the response of
single crystals of hard brittle materials to shock-wave loading. In the experiment
with ruby, the peak stress did not exceed the Hugoniot elastic limit (HEL). The
interface velocity history is smooth and mimics the shape of the stress pulse inside
the sample. The high velocity pullback indicates the dynamic tensile strength (the
spall strength), is as high as 10 GPa. In the other shot, the peak stress exceeded the
HEL. The spall strength drops practically to zero, and irregular oscillations appear
in the wave profile. Vanishing resistance to tension after shock compression
above the HEL was observed for quartz single crystals as well.17
Presumably, the
absence of fracture nucleation sites enables high spall strength at peak shock
stresses below the HEL. However, fracture nucleation sites obviously appeared
during shock compression above the elastic limit.
The high-frequency particle velocity jitter is evidence for heterogeneity of the
inelastic deformation process. Similar records have been obtainable for quartz17
and olivine18
. Another characteristic feature is the significant stress relaxation
behind the elastic precursor front caused by intense multiplication of the
deformation carriers. This also is typical for brittle crystals.18,19
Figure 3 presents the results of measurements of shock compressibility of
sapphire crystals. Above the HEL, the Hugoniot shows a collapse toward the
isotropic compression curve (hydrostat): the stress offset above the HEL is 3.8 to
4.3 GPa, while at the HEL it ranges from 5.5 to 11 GPa. The collapse of the
Hugoniot indicates a collapse of shear strength. Quartz22
, magnesium oxide19,23
,
zirconia24
, iron-silicate almandine-garnet25
, and olivine single crystals18
also
The HEL is limiting stress for linear elastic compression under uniaxial strain;
e.g., it is the compressive strength for full lateral confinement.
200 Ceramic Armor Materials by Design
exhibit this phenomenon. Grady19
found that in magnesium oxide crystals,
strength persists at the Hugoniot state and the release paths deviate significantly
from the hydrostat; the initial elastic release velocities was some below expected
longitudinal (elastic) velocities but was substantially above expected bulk
(inelastic) sound velocities.
0.0 0.2 0.4 0.6
0.0
0.2
0.4
0.6
0.8
1.0
Expected
spall signal
Sapphire
Ruby
Ve
loc
ity
, k
m/s
Time, s
0,85 0,90 0,95 1,00
0
15
30
45
K0=226 GPa,
K'=4.0
Str
ess, G
Pa
V/V0
Figure 2. The wave profiles at interface between alumina single crystals and water
window at shock wave loading up to various peak stresses.
Figure 3. Stress-volume relations for sapphire under shock-wave compression.
Points present the data.20,21
The dot-dashed line shows the isotropic compression
curve.21
Wang and Mikkola examined recovered sapphire samples with transmission
electron microscopy after shock compression up to 23 GPa.26
They observed a
significant number of slip bands in different crystallographic directions and
suggested that a large amount of plastic deformation had occurred at shock
stresses of 12 GPa and more. Anan’in et al. have revealed glass-like interlayers
between quartz blocks in recovered single crystals after shock loading.27
This
lamellae structure indicates a heterogeneous nature of shock deformation of
quartz, accompanied by melting. Grady has developed a model of localized
dissipation of elastic strain energy in low thermal conductivity strong solids.28
The dissipation leads to local temperature growth, which reduces the local flow
stress, causing the shear strain and the energy release to be localized within
narrow bands where the temperature may reach melting. However, no signs of
local melting were observed in sapphire and other hard materials.
Thus, hard single crystals show a more or less substantial reduction in shear
strength at shock compression beyond their Hugoniot elastic limits. Within the
elastic range, they demonstrate very high dynamic tensile strength, which is
Ceramic Armor Materials by Design 201
attributed to lack of flaws and heterogeneities. At shock compression above the
HEL, they show vanishing dynamic tensile strength.
Unlike crystals, silicate glasses maintain high tensile strength after shock
compression above the HEL. Figure 4 presents the free surface velocity profiles
for K8 crown glass.29
Spallations were not observed in these shots, which means
that the spall strength of the glass exceeds 6.8 GPa when the shock stress is below
the HEL, and it remains very high above the HEL. For comparison, the static
tensile strength of glasses is around 0.1 GPa. The reason for such a large
discrepancy is that the fracture nucleation sites in homogeneous glass are
concentrated on the surface. These incident microcracks are activated and
determine the strength magnitude in the static measurements, whereas spall
strength is an intrinsic property of matter.
0,0 0,5 1,0 1,5 2,0
0,0
0,6
1,2
1,8
Reflection from at a failed layer
1Simulation
2
Calculated rereflectionat the impact surface
Fre
e S
urf
ace
Ve
locity (
km
/s)
Time ( s)
Figure 4. Experimental results for K8 glass samples 6.1 mm thick at the impact
velocities of 670 30 m/s (profile 1), and 1900 50 m/s (profile 2). Impactors are
steel 0.9 mm thick (1) and aluminum 2 mm thick backed by paraffin (2). Dashed
line shows results of computer simulations assuming no failure.
At high pressures, brittle glasses become ductile. Ductility of glasses is
possible because there is a loose microstructure with a large amount of molecular-
size voids. It is known that glasses show gradual structural changes, resulting in
increased density.30
Irreversible densification of some glasses also occurs under
shock compression above the HEL.31,32
It is supposed that the irreversible
densification and compaction in the silicate structure are responsible for the
plastic flow properties of glasses under high pressure.33
Once the plastic flow
starts, stress relaxation reduces the stress concentration at crack tips and thus stops
the propagation. The high spall strength revealed in the stress range above the
HEL means that the ductility is preserved during the subsequent tensile loading.
202 Ceramic Armor Materials by Design
Single crystals and glasses are initially homogeneous in bulk so the only way
fracture nucleation sites may be formed is in the course of plastic deformation. In
this sense the difference between single crystals and glasses is that hard crystals
have only a limited number of crystal planes and directions in which the usual
mechanisms of ductility may work, whereas the ductility of amorphous glasses is
completely isotropic. The impossibility of plastic shear along arbitrary directions
in crystals results in stress concentration at points of intersections of slip bands or
twins that, in turn, may result in cracking at compression or unloading.
Comparison of the measured free surface velocity history (curve 1 in Fig. 4)
with results of computer simulations shows that the elastic wave reverberation
inside the glass plate sample occurs earlier than expected. The early arrival of
second compression pulse is due to reflection from a failure wave at some
distance from the impact surface.
The failure wave is a network of cracks that are nucleated on the surface and
propagate into the stressed body. There are many observations of fracture front
propagation in glasses under tensile stresses. Schardin recorded expansion of
fractured areas with a sharp front formed by bifurcated cracks.34
Galin et al.
reported an explosion-like fracture under bending of high-strength glass with
removed surface defects.35
The explosion-like fracture was treated by Galin and
Cherepanov, who termed it a self-propagating failure wave.36
The similar fracture mode under compression was revealed in shock-wave
experiments with glass plates. Some results of observations of failure wave
phenomena were reviewed recently.29,37
The observations may be summarized as
follows: (i) failure waves are observed when the impact stress exceeds some
threshold but is still below the Hugoniot elastic limit of glass; (ii) failure waves
nucleate at a plate surface; (iii) decrease of deviator stresses and vanishing of
tensile strength occur behind the failure wave front; (iv) propagation of the failure
wave stops when the stress in front of it decreases; and (v) the failure wave
velocity is much less than the sound speed. Many measurements give the failure
wave speed equal to an ultimate speed of growth of cracks (~1.5 km/s for glass),
but higher and lower velocity values were reported as well. Both constant and
decreasing propagation velocities were reported.
Failure waves present a mode of catastrophic fracture in an elastically
compressed media that is not limited to impact events. One hopes that the
investigations of failure waves in shock-compressed glass will provide
information about the mechanisms and general rules of nucleation, growth, and
interactions of multiple cracks and lead to better understanding of experiments
with other hard brittle materials, such as ceramics and rocks.
DYNAMIC STRENGTH PROPERTIES OF POLYCRYSTALINE CERAMICS
Modern shock-wave tests of ceramics include measurements of the Hugoniot
over a wide stress range, shock front rise time, Hugoniot elastic limit (HEL),
Ceramic Armor Materials by Design 203
stress state immediately after shock compression, high pressure stress strain path,
tensile (spall) strength after shock compression below and above the elastic limit,
and post-test examination of recovered samples. Fig. 5 shows typical particle
velocity histories for SiC and B4C ceramics.38
The wave profiles in Fig. 5 exhibit two extreme examples of behavior of
ceramics in plane shock waves. The response of silicon carbide is very similar to
that of ductile materials. There is an initial elastic arrival whose amplitude (the
HEL) is the limit stress for elastic behavior, there is a second shock corresponding
to bulk compression, followed by elastic and bulk unloading waves that originate
from the rear surface of the flyer plate. Post-yield strength of silicon carbide,
determined by comparison of uniaxial strain and calculated hydrodynamic
response, increases considerably beyond the initial dynamic yield. The release
trajectories for silicon carbide indicate reverse yielding (e.g. reversal of the sense
of shear) and continued elastic-plastic bulk behavior, probably with a Baushinger
effect at higher peak stresses. The shock response of B4C is quite different. The
bulk compression wave is much slower. According to Grady, the Hugoniot
collapses to the hydrostat at stresses approaching about twice the HEL.39
A
dispersed character of the unloading wave indicates inelastic strain starts almost
immediately behind the rarefaction wave front. The stress-strain trajectory for the
B4C ceramic 38,39
shows evidence of dilatancy when the compressive stress
approaches zero on unloading.
Correspondingly, post-yield characteristics of the materials are qualitatively
contrasted by the shape of the bulk compression waves. For silicon carbide,
positive slope of the wave demonstrates strain hardening. The stress drop after the
HEL in boron carbide, in contrast, indicates post-yield softening. Spall strength is
sustained for shocks above the HEL in SiC, but not in B4C.
0,0 0,5 1,0 1,5
0,0
0,5
1,0
1,5
SiC
B4C
Pa
rtic
le V
elo
city, km
/s
Time, s
Figure 5. Particle velocity profiles for SiC and B4C ceramics measured at the
interface with a LiF window.38
204 Ceramic Armor Materials by Design
The dynamic compressive properties of many ceramics have been summarized
recently in our report.44
Here we’ll consider, as an example, the published data
for alumina. Table I and Figs. 6 and 7 present the shock data (longitudinal sound
speed, cl, the Hugoniot elastic limit, HEL, and the Von Mises yield stress,
Y = HEL (1-2 ) / (1- ) as functions of the initial density, 0) for different Al2O3
ceramics. Whereas the general trend is HEL reduction with increasing ceramic
porosity, the impurity content, the grain size, and material processing also
influence the HEL value. The compaction of more porous ceramics occurs within
the stress range from the yield point to about 30 GPa. At the higher stresses the
states of all alumina ceramics are practically described by one curve in stress-
volume coordinates. Beyond the compaction region, the yield strength, Y,
estimated from the stress offset between the Hugoniot and isotropic compression
curve is comparable to but somewhat smaller than the yield strength at the HEL.
The profiles of shock compression waves propagating through alumina
ceramics exhibit an elastic jump and a subsequent dispersed rise to the bulk wave
which compresses the matter to a final state. This gradual transition from the
elastic to inelastic portions of the compression wave is typical for strain hardening
materials. Cagnoux and Longy measured the free surface velocity profiles for
alumina at various rise times of the compression wave entering into the sample.45
The HEL was found to be independent of the wave propagation distance, the peak
shock stress, and the entering stress gradient; this means that there is no influence
of strain rate on the yield strength of alumina in a range of 5 104 to 6 10
5 s
-1. On
the other hand, Furnish and Chhabildas found evidence of rate-dependent
behavior of AD995 ceramic at step-like compression.46
According to many
measurements, the unloading wave front in shock-compressed alumina is elastic;
however, there is no sharp distinction between the elastic and inelastic parts of
unloading.
Longy and Cagnoux found that alumina ceramics with 2% porosity but with
grain size of 5 or 60 m exhibit a HEL of 8.5 GPa and 5 GPa, respectively.56
Microscopic examination of impure alumina showed microcracks in the inter-
granular glassy phase after shock stress 0.9 HEL, and there was no correlation
between the HEL and microcrack levels.
Cagnoux carried out microscopic examination of alumina samples of two
different grain sizes (4.7 m and 10-20 m) with 99.7% Al2O3 content and 3.91
g/cm3 density.
52The samples were recovered after compression above their HEL
by spherical shock waves. In the region of maximum peak stresses, the fine-grain
alumina remained uncracked, whereas the coarse-grain sample was
microfragmented. SEM photographs of fine-grain samples showed reduction in
porosity, with no slip-nucleated microcracks; in the coarse-grained sample,
numerous twins were observed. It was concluded that twinning is favored by large
grain size, while slipping by small grain size.
Ceramic Armor Materials by Design 205
Table I. Hugoniot elastic limits of aluminas
Material (wt. fractionAl2O3), Grain Size
0, g/cm3,
(Porosity, %)
cl, km/s Poisson’s
ratio,
HEL,GPa
Y, GPa Ref.
Lucalox (99.8%) 3.98 (<0.2) 10.95 0.2363 11.2 1.3 7.8 43
Lucalox (99.9%),
25-40 m
3.969 10.92 9.1 0.4 6.0 44
MTU JS-I (99.99%), 1.5
m
3.974 10.9 0.237 11-11.9 7.6-8.2 45
D999 (99.9%), 4 m 3.99 10.82 0.232 13-14 9-9.8 46
Carborundum hot pressed 3.92 (0.8) 10.59 0.243 9.2-16 47
Wesgo Al-995 (99.5%) 3.81 (3.5-4.3) 10.2 0.218 8.3 0.5 6.0 43
D975 (97.5%), 4 m 3.8 10.3 0.234 7.5-9 5.2-6.2 46
Coors AD995,aluminosilicate glassbinder
3.88 (2) 10.56 6.7 0.1 48
Coors AD995 3.89 10.59 0.234 6.2 0.4 4.3 49
Coors AD-85 3.42 (6.6) 8.84 0.256 6.1-6.5 4.1 47
Coors AD-85 (84%) 3.42 (6.6) 4.7-6.1 50
H880 (88%), 2 m 3.55 9.1 0.226 5.5-6.5 3.9-4.6 46
Diamonite P-3142-1 3.72 (5.5) 9.98 0.234 7.2-8.1 47
Desmarquest alumina 3.62 (5.3) 9.45 4.5 45
ENSCI, 4.7 m 3.91 (2) 10.63 8.7 0.4 51
ENSCI, 1 m 3.54 (11) 9.34 5 51
ENSCI, 0.6 m 3.31 (17) 8.55 4-5 51
ENSCI T60 (99.7%), 5-
125 m
3.85 (3.5) 10.32 5 51
UL500 (93.8%), 11 m 3.62 (6.2) 9.77 6.5 51
3.2 3.4 3.6 3.8 4.0
8.5
9.0
9.5
10.0
10.5
11.0
cl, k
m/s
Density, g/cm3
Figure 6. Longitudinal sound speed, cl, in different Al2O3 ceramics as a function
of their density.
206 Ceramic Armor Materials by Design
3.2 3.4 3.6 3.8 4.00
3
6
9
12
15
Lucalox
5-125 m
AD995AD-85
0.5-5 m
HE
L,
GP
aDensity, g/cm
3
Figure 7. Hugoniot elastic limits of Al2O3 ceramics of different density and grain
size.
Table II summarizes the results of rod impact experiments which provide the
dynamic failure threshold under uniaxial stress conditions.76
Comparison of data
of Tables I and II shows that the failure threshold for rod impacts is lower than the
dynamic stress Y at the HEL for the same material. For the ceramic rods mounted
within a close-fitting high-impedance sleeve, the maximum axial stress may
approach the HEL value.
Table II. The failure thresholds for alumina ceramics in rod impact experiments.
Material, impact
conditions
Confine-
ment
Failure Threshold, GPa Reference
AD94, direct impact - 2.7 54
AD99, direct impact - 4 54
AD995, direct impact - 3.15 55
AD995, direct impact Ta 5.8 6.3 55
AD995, dispersed impact - 3.5 4.2 53
AD995, dispersed impact Steel 4.6 53
AD995, direct impact - 3.6 3.7 56
AD995, direct impact Steel 4.2 56
AD995, direct impact - 3.8 57
Figure 8 presents results of spall strength measurements as a function of
normalized peak stress for alumina ceramics.45,48,58,59
It seems the spall strength
undergoes a transition, first decreasing near the HEL, then increasing with
increasing pressure above the HEL. The reduction in spall strength value near the
HEL is especially significant for aluminas with a large glassy phase content. This
observation correlates with observed 51
microcracks in the inter-granular glassy
Ceramic Armor Materials by Design 207
phase at shock stress of 0.9 HEL and over. Within the elastic region, the spall
strength decreases with increasing porosity and grain size.
0,1 1
0,0
0,5
1,0
1,5
0,3 0,5 432
Pure hot pressed,
HEL=10-12 GPa
MTU JS-1, HEL=11.9 GPa
AD995, HEL=6.7 GPa
AD85, HEL=6GPa
Spa
ll S
tre
ngth
, G
Pa
(Peak Stress) / HEL
Figure 8. Spall strength of alumina ceramics as a function of peak stress. Data for
AD8558
, for and for hot pressed pure alumina.45,59
BRITTLE FAILURE CRITERIA AND MODELS
Reviews of models and criteria of brittle fracture in quasi-static compression
are available.3,60
Most models consider failure as a consequence of three
sequential events: crack initiation, crack propagation, and crack coalescence.
According to the Griffith’s criterion, crack initiation occurs when the highest local
tensile stress at the longest crack of the most dangerous orientation reaches a fixed
critical value. For a biaxial stress state, the corresponding relationship is
( 1- 2)2
8 f ( 1+ 2)=0, (1)
where 1, 2 are principal stresses, f is a material constant, which is assumed to
be the ordinary tensile stress for uniaxial stressing. Thus, Griffith’s criterion
predicts the value of the uniaxial compressive strength to be eight times the value
of the uniaxial tensile strength. This ratio is smaller than the ratio commonly
measured for rocks and other brittle materials.
Chen and Ravichandran have found that Mohr-Coulomb-like criterion,
generally used as a yielding criterion for sand-like materials, fits rather well the
experimental data on failure of ceramics under lateral confinement.61,62
The
criterion is expressed by a relationship
208 Ceramic Armor Materials by Design
+ p = 0, (2)
where is the absolute value of shear stress, p is the pressure or the mean
stress, 0 is the shear strength of the material without any pressure, and is the
internal friction coefficient. It is expected that the value is between -1.5 and 0.
For AlN ceramic, = -1; for Macor glass ceramic, = -0.74. Actually, in the
Coulomb’s interpretation 63
, there should be normal stress component
perpendicular to the shear direction used in Eq. (2) instead of the mean stress. In
this case, the competition between the resolved shear force and the friction
forces p explains the failure angle <45 observed in rocks and granular materials
under compression.
In general, theoretical criteria, such as that of Griffith, often give an
inadequate fit to the data. Because of that, empirical criteria have been developed
to meet the practical requirements of accurate strength prediction and simplicity of
use. Since different mechanisms of inelastic deformation and failure may operate
depending on the region of stress space, different strength criteria should be used
at different levels of stress. For computations of impact problems, continuum
damage mechanics models are needed. Such models operate with defined damage
parameters and include an evolutionary equation for the damage and a constitutive
equation that relates the stress and strain to the damage. In this sense the model of
Ashby and Sammis is very illustrative (Figure 9).64
Yield
YieldFailure
Damage initiation
Failure
Damage initiation
-1
-2=-
3
Figure 9. Failure and yield surface according to Ashby and Sammis.64
COMPARISON OF 1-D-STRESS AND 1-D-STRAIN DATA FOR CERAMICS
AT VARIOUS STRAIN RATES
To what extent do the strength properties of ceramics depend on strain rate?
For many metals the strain rate sensitivity of the flow stress increases steeply
above ~103 – 10
4 s
-1. This is interpreted as a transition in the rate-controlling
Ceramic Armor Materials by Design 209
mechanism of dislocation motion. For low rates the motion is aided by thermal
fluctuations; at very high strain rates, viscous phonon drag becomes dominant.
For brittle materials, cracking accompanies the inelastic deformation, so the
kinetics of cracks nucleation and growth may contribute into a total strain-rate
dependency of the resistance to inelastic deformation.
Conclusions about rate sensitivity of ceramics should be based on comparison
of compressive strength properties over a wide strain rate range. However, this
procedure is not easy. In Fig. 10 the yield strength, Y, determined from the HEL
by the Von Mises relationship, is compared with the failure stress values
measured in uniaxial stress conditions (quasistatic tests, Hopkinson bars, rod
impact) for alumina. The data indicate weak rate dependencies at strain rates less
< 103 s
-1 and > 10
5 s
-1 with a sharp transition between these ranges.
10-6
10-3
100
103
106
0
3
6
9
12
Rod impact
Y/(1-2 )
Y
Uniaxial Stress
Un
iax
ial s
re
ain
(S
ho
ck
)
A lumina
Fa
ilu
re
Str
es
s,
GP
a
Strain Rate, s-1
Figure 10. Dynamic failure properties of alumina ceramic under quasi-static and
dynamic compressive loading. Solid points reproduce data65
for failure strength
measured under uniaxial stress conditions and the yield strength at HEL (1D
strain) calculated with Von Mises yielding criterion. Hatched rectangle shows the
region of rod impact data. Open points are the failure strength values multiplied
by a factor of 1/(1-2 ) according to Rosenberg.66,67
Rosenberg, however, has suggested the use of Griffith’s failure criterion for
reconciliation 1D stress and 1D strain data instead of yield criteria.66,67
In this
case the HEL should be
HEL = (1- )Yc / (1-2 )2, (3)
where Yc is the compressive strength under 1D stress conditions. Thus, the
Griffith’s failure criterion predicts the HEL strength to be higher by a factor of
1/(1-2 ) than the yield criteria used for ductile materials, and this reduces the
discrepancy between the 1D stress and 1D strain data. Figure 10 shows that the
210 Ceramic Armor Materials by Design
Griffith’s failure criterion provides very good agreement between 1D stress and
1D strain dynamic data.
While the rod impact and Hopkinson bar experiments permit a fairly confident
judgment that failure is brittle, interpretation of the shock-wave experiments
under 1-D strain conditions is more ambiguous. The Hugoniot data and the shock-
wave profiles are not by themselves sufficient to make definite conclusions about
ductile or brittle response. The continuity in spall strength values indicates that, in
many cases, the mode of deformation is mainly ductile.
DISCUSSION
It seems the compressive fracture is (or may be) a secondary effect of the onset
of plastic deformation in brittle materials. Experiments with single crystals
demonstrate that microcracks may nucleate at stress concentrators formed under
inelastic compression, even if initially the material does not contain any defects.
The comparison of the behaviors of single crystals and glasses shows that the
choice of ductile or brittle response is controlled by the capability of material to
accommodate local shears in different directions.
Most of the shock-wave tests of polycrystalline ceramics, with the exception
of boron carbide, do not show unambiguous evidences of fracture under uniaxial
compression. Only boron carbide manifests dilatancy effects on the stress-strain
trajectory when the compressive stress approaches zero on unloading. Even the
observed decrease in spall strength after shock compression above the HEL may
be the result of cracking not under compression but at the end phase of unloading
as an effect of transversal stresses when the axial stress becomes zero. Additional
efforts are needed to answer the question of whether or not microcracking of
ceramic materials occurs under shock compression.
Shock wave tests provide information about the resistance to inelastic
compressive deformation, which is very useful whether or not microcracking
occurs. Additional information over a wider range of stressed states and strains
may provide experiments with spherical 68
and cylindrical 69
divergent shock
waves and experiments with granular ceramics and ceramic powders 70
.
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using axi-symmetric shock waves”; pp. 489-492 in Shock Compression of Condensed Matter – 1997. Edited by S.C. Schmidt, D.D. Dandekar, and J.W.
Forbes. AIP Conference Proceedings 429, 1998. 70
L.W. Meyer and I. Faber, “Investigations on granular ceramics and ceramic
powder,” J. Phys. IV France 7 [C3] 565-570 (1997).
ACKNOWLEDGEMENT The research reported in this document was performed in connection with Contract number DAAD17-
01-D-0001 with the U.S. Army Research Laboratory. The views and conclusions contained in this document are those of the authors and should not be interpreted as presenting the official policies or position, either expressed or implied, of the U.S. Army Research Laboratory or the U.S. Government unless so designated by other authorized documents. Citation of manufacturer’s or trade names does not constitute an official endorsement or approval of the use thereof. The U.S. Government is authorized to reproduce and distribute reprints for Government purposes notwithstanding any copyright notation hereon.
216 Ceramic Armor Materials by Design
RECENT DEVELOPMENTS IN SPLIT HOPKINSON PRESSURE BAR
TESTING
W. Chen, and B. Song D. J. Frew and M. J. Forrestal
The University of Arizona Sandia National Laboratories
Tucson, AZ 85721-0119 Albuquerque, NM 87185-1174
ABSTRACT
The split Hopkinson pressure bar (SHPB) technique has been widely used to
determine the dynamic properties of engineering materials. Recent applications
of this technique to materials with extreme properties (e.g., extremely soft or
extremely hard) have forced more careful examination of the SHPB apparatus and
mandated modifications to this well-established experimental technique in order
to obtain valid results. In this paper, we present a brief review of recent
developments that adapt this technique for testing a variety of engineering
materials under valid dynamic testing conditions. It is shown that, in order to
subject the specimen to uniform deformation at a nearly constant strain rate under
dynamic stress equilibrium, pulse shaping must be used in the SHPB experiments.
In addition, a more sensitive transmission bar must be employed to detect weak
transmitted signals when testing soft materials.
INTRODUCTION
The SHPB or Kolsky bar technique originally developed by Kolsky1, 2
has
been used by many investigators to obtain dynamic compressive properties of
solid materials. The evolution of this experimental method and recent advances
are discussed by Nicholas3, Ellwood et al.
4, Franz et al.
5, Follansbee
6, Nemat-
Nasser et al.7, Ramesh and Narasimhan
8, Gray
9, and Gray and Blumenthal.
10 This
technique has mostly been used to study the plastic flow behavior of metals that
undergo large strains at strain rates between 102 – 10
4/s. As discussed by Yadav
et al.11
, data for the compressive flow stress of metals are typically obtained for
strains larger than a few percent because the technique is not capable of
measuring the elastic and early yield behavior. Also, the method of measuring
strains with resistance strain gages mounted on the metal bar surfaces limits the
applicability of SHPB to test the dynamic properties of low-strength and low-
impedance materials. Recently, there has been a need to obtain reliable dynamic
Ceramic Armor Materials by Design 217
To the extent authorized under the laws of the United States of America, all copyright interests in this publication are the propertyof The American Ceramic Society. Any duplication, reproduction, or republication of this publication or any part thereof, withoutthe express written consent of The American Ceramic Society or fee paid to the Copyright Clearance Center, is prohibited.
material properties for brittle materials with failure strains are less than about 1.0
percent, such as ceramics12,13
and rocks14
, and for soft material that are too week
to generate a clear transmitted pulse, such as silicone rubbers15
and polymeric
foams16
. These brittle and soft materials require careful examination of the SHPB
technique and mandate necessary modifications to obtain valid data. In this
paper, we briefly review the SHPB technique, the challenges it faces, and the
remedies for obtaining valid experimental results.
As shown in Fig. 1, a conventional SHPB consists of a striker bar, an incident
bar, a transmission bar, and a sample placed between the incident and
transmission bars. Nicholas3, and Gray
9 present the working principles and the
equations that describe the sample response in terms of the measured strain
signals.
Striker
BarIncident Bar Transmission Bar
Specimen
( s, cs, As)
i r
( , c, A) ( , c, A)
2
t
1
lo
Fig. 1 A schematic illustration of a split Hopkinson pressure bar.
The equations that relate the strain gage signals to material responses
(Nicholas3, and Gray
9) are based on idealized 1-D wave propagation analysis. In
addition, it is assumed that the sample undergoes homogeneous deformation and
is in dynamic stress equilibrium. In a conventional SHPB experiment, the wave
propagation in the elastic bars is not as ideal as assumed. Figure 2 presents a set
of oscilloscope record of the incident, reflected and transmitted strain signal
during a typical SHPB experiment on a 1046 mild steel specimen. The signals in
Fig. 2 show that the incident pulse is not rectangular in shape as idealized in 1-D
wave theory. The reflected signal, which directly correlates to the strain rate in
the specimen, has significant fluctuations with its average amplitude decreasing
with time. This indicates that the specimen is not deforming at a nearly constant
strain rate. Furthermore, due to the significant fluctuations in the beginning
portion of the reflected pulse, the initial specimen strain from a SHPB test is not
considered to be reliable, leading to an unreliable Young's modulus of a material
measured with a SHPB. The dynamic stress state in the specimen can be
examined using a 1-wave, 2-wave method10, 17, 18
, which compares the axial force
history on the transmission side of the specimen (transmitted signal) to that on the
218 Ceramic Armor Materials by Design
incident side (the difference between the incident and reflected signals). Figure 3
shows the results of such a 1-, 2-wave analysis for the experiment in Fig. 2.
Fig. 2 Oscilloscope record of a conventional SHPB experiment on a mild steel.
Fig. 3 Axial force histories on the specimen by 1-wave, 2-wave analysis.
The results shown in Fig. 3 illustrate that dynamic equilibrium in the
specimen is approximately reached during the later stages of the experiment (t >
~60 s). During the early stages of the experiment (t < ~30 s), the specimen is
not in stress equilibrium. This leads to uncertainties on the dynamic yield
strength as determined by SHPB tests. Gray and Blumenthal10
reached similar
conclusions regarding the equilibrium of a soft rubber specimen.
Ceramic Armor Materials by Design 219
When the SHPB is used to test relatively brittle materials such as ceramics12, 13
and rocks14
, most of the material behavior of interest occurs at strains less than
about 1.0 percent, which is within the large error range for conventional SHPB
experiments. For other material such as polymers and shape-memory alloys,
loading history significantly affects the mechanical behavior. The initial portion
of the experimental duration cannot be ignored. Therefore, to use SHPB
technique for obtaining the dynamic properties of such materials, modifications
must be made to the testing technique to ensure that the specimen deforms
uniformly at a nearly constant strain rate under dynamically equilibrated stress.
Such modifications include pulse shaping, a sensitive transmission bar, and
dynamic equilibrium monitoring. Next, we briefly describe each of these new
developments in SHPB testing in which we have been involved.
PULSE SHAPING
The initial significant fluctuations in the reflected signal shown in Fig. 2 and
the initial non-equilibrium shown in Fig. 3 indicate that the incident loading pulse
profile needs to be controlled to facilitate stress equilibrium and uniform
deformation in the specimen. Some of the advantages and necessities for shaping
the incident pulse for SHPB experiments were discussed twenty years ago. Franz
et al.5 and Follansbee
6 wrote review papers that discussed pulse shaping for SHPB
experiments with metal samples. In these review papers, the authors emphasized
that a slowly rising incident pulse is preferred to a pulse that rises steeply in order
to minimize the effects of dispersion and allow the sample to achieve dynamic
stress equilibrium. Franz et al.5 and Follansbee
6 also discuss experimental
techniques for pulse shaping and a numerical procedure17
for correcting raw data
for wave dispersion in the bars. These authors5 and Ellwood et al.
4 show that a
properly chosen pulse shaper can also be used to generate a nearly constant strain
rate in the sample. Gray9 and Gray and Blumenthal
10 present additional
information in recent survey papers that include these subjects. However, Duffy
et al.19
were probably the first authors to use pulse shapers to smooth pulses
generated by explosive loading for the torsional Hopkinson pressure bar.
While pulse shaping techniques have been successfully used to achieve the
goals of many different experiments, pulse shapers are usually designed by
experimental trials that exclude a model to guide the design parameters. For
examples, Wu and Gorham18
used paper on the impact surface of the incident bar
to eliminate high frequency oscillations in the incident pulse for Kolsky
compression bar experiments. Togami et al.20
used a thin, Plexiglas disk to
produce nondispersive compression pulses in an incident bar, and Chen et al.21
used a polymer disk to spread the incident compressive pulses for experiments
with elastomers. Christensen et al.22
used striker bars with a truncated-cone on
the impact end in an attempt to produce ramp pulses. In contrast to other pulse
220 Ceramic Armor Materials by Design
shaping studies, Nemat-Nasser et al.7 modeled the plastic deformation of an
OFHC copper pulse shaper, predict the incident strain pulse, and show good
agreement with some measured incident strain pulses. Frew et al.23
further
extended the model to describe the behavior of C11000 copper pulse shapers
driven to much larger strains. In the equations that govern wave propagation in
the striker bar, it was found that the added mass from the sabot must also be
considered24
.
Pulse-shaping techniques have been applied recently to obtain valid dynamic
material properties of a variety of materials, such as limestone14
, ceramics23
, a
shape-memory alloy25
, a polymeric foam16
, and rubbers15
. Pulse shaping has also
been used in dynamic tension experiments26
. As an example, Fig. 4 shows the
oscilloscope record of a pulse-shaped SHPB experiment on the same 1046 mild
steel. The incident pulse was created by placing a combination of hardened 1046
steel and C11000 copper disks on the striking end of the incident bar. The nearly
flat reflected signal shows minimum fluctuations, which indicates that a nearly
constant strain rate has been achieved in the specimen. Furthermore, a detailed
examination of the reflected signal reveals that, without the fluctuations
associated with the reflected signal (Fig. 2), the reflected signal in Fig. 4 actually
is composed of two plateaus: a small-amplitude plateau followed by a second,
larger one. Data reduction further reveals that the small plateau corresponds to
the elastic deformation of the specimen, whereas the larger one is associated with
the plastic flow. In addition to revealing the details in the reflected pulse, pulse
shaping also facilitates dynamic equilibrium. Figure 5 shows the 1-, 2-wave
analysis for the pulse shaped experiment, which indicates a nearly perfect
agreement between the front- and back-end force histories.
SENSITIVE TRANSMISSION BAR
When a soft material is tested with a SHPB, the transmitted signal can be too
weak to provide a stress history for the specimen21
. More sensitive transmission
bars are thus necessary. Low-impedance bars, such as polymer bars, will extend
the time for the sample to reach dynamic equilibrium14
. We have developed an
aluminum transmission tube21
and a quartz-crystal embedded aluminum bar27
to
provide high sensitivity of the transmission bar, while still maintaining the high
impedance mismatch between the specimen and the bar.
DYNAMIC EQILIBRIUM MONITORING
When the specimen is a very soft material, nearly all of incident signal is
reflected back into the incident bar. This introduces significant errors in the 2-
wave analysis, which takes the differences between the incident and reflected
signals. We have developed quartz-crystal methods to directly measure the front-
Ceramic Armor Materials by Design 221
and back-end force histories in the specimen15, 16
, which directly monitors the
dynamic equilibrium process in the specimen.
Fig. 4 Oscilloscope record of a pulse-shaped SHPB experiment on a mild steel.
Fig. 5 Axial force histories after pulse shaping by 1-wave, 2-wave analysis.
SUMMARY
Pulse shaping must often be employed to obtain dynamic material properties
with a SHPB to ensure that the specimen is deforming uniformly at a nearly
constant strain rate under dynamic equilibrium. Proper modifications to a
conventional SHPB are necessary when testing hard or soft materials.
222 Ceramic Armor Materials by Design
ACKNOWLEDGEMENTS
This work was sponsored by the U.S. Army Research Office through a grant
to The University of Arizona (G-DAAD19-00-1-0493) and the Sandia National
Laboratories Joint DoD/DOE Penetration Technology Program. Sandia is a
multi-program laboratory operated by Sandia Corporation, a Lockheed Martin
Company, for the U.S. Department of Energy under Contract DE-AC04-
94AL8500.
REFERENCES 1H. Kolsky, “An Investigation of the Mechanical Properties of Materials at
Very High Rates of Loading,” Proc. Royal Soc. Lond., B62 676-700, (1949).2H. Kolsky, Stress Waves in Solids. Dover, New York (1963).
3T. Nicholas, “ Material Behavior at High Strain Rates,” Impact Dynamics,
Chapter 8, John Wiley & Sons, New York, (1982). 4S. Ellwood, L. J. Griffiths, and D. J. Parry, “Materials Testing at High
Constant Strain Rates,” J. Phys. E: Sci. Instrum., 15 280-282 (1982).5C. E. Franz, P. S. Follansbee, and W. J. Wright, “New Experimental
Techniques with the Split Hopkinson Pressure Bar,” in the 8th
Int. Conf. on High
Energy Rate Fabrication, ASME, (ed. I. Berman and J. W. Schroeder), San
Antonio, TX, June 17-21 (1984). 6P. S. Follansbee, “The Hopkinson Bar,” Mechanical Testing, Metals
Handbook, 9th
ed., 8, Am. Soc. for Metals, Metals Park, Ohio, 198-217 (1985). 7S. Nemat-Nasser, J. B. Isaacs, and J. E Starrett, “Hopkinson Techniques for
Dynamic Recovery Experiments,” Proc. R. Soc. Lond., A435 371-391,(1991). 8K. T. Ramesh and S. Narasimhan, “Finite Deformations and the Dynamic
Measurement of Radial Strains in Compression Kolsky Bar Experiments,” Int. J.
Solids Structures, 33 3723-3738 (1996). 9G. T. Gray, “Classic Split-Hopkinson Pressure Bar Technique,” ASM
Handbook, 8, Mechanical Testing and Evaluation, ASM International, Materials
Park, OH, 44073-0002 (2000). 10
G. T. Gray and W. R. Blumenthal, “Split-Hopkinson Pressure Bar Testing of
Soft Materials,” ASM Handbook, 8, Mechanical Testing and Evaluation, ASM
International, Materials Park, OH, 44073-0002 (2000).
11S. Yadav, D. R. Chichili, and K. T. Ramesh, “The Mechanical Response of
a 6061-T6 Al/Al2O3 Metal Matrix Composite at High Rates of Deformation,”
Acta metall. Mater., 43 4453-4464, (1995). 12
W. P. Rogers and S. Nemat-Nasser, “Transformation Plasticity at High
Strain Rate in Magnesia-Partially-Stabilized Zirconia,” J. Am. Ceram. Soc., 73
136-139, (1990).
Ceramic Armor Materials by Design 223
13W Chen. and G. Ravichandran, “Dynamic Compressive Failure of a Glass
Ceramic Under Lateral Confinement,” J. Mech. Phys. Solids. 45 1303-1328,
(1997).14
D. J. Frew, M. J. Forrestal, and W. Chen, “A Split Hopkinson Bar
Technique to Determine Compressive Stress-Strain Data for Rock Materials,”
Experimental Mechanics, 41 40-46, (2001).15
W. Chen, F. Lu, D. J. Frew, and M. J. Forrestal,“Dynamic Compression
Testing of Soft Materials,” ASME J. Appl. Mech., rivised (2001). 16
W. Chen, F. Lu, and N. Winfree, "High Strain-Rate Compressive Behavior
of a Rigid Polyurethane Foam with Various Densities," Experimental Mechanics,
accepted (2001).17
P. S. Follansbee and C. E. Franz, “Wave Propagation in the Split Hopkinson
Pressure Bar,” Trans. ASME, J. Eng. Mat. Technol., 105 61-66 (1983).18
X. J. Wu and D. A. Gorham, “Stress Equilibrium in the Split Hopkinson
Pressure Bar Test,” J. Phys. IV France, 7 (C3) 91-96 (1997).19
J. Duffy, J. D. Campbell, and R. H. Hawley, “On the Use of a Torsional
Split Hopkinson Bar to Study Rate Effects in 1100-0 Aluminum,” ASME J. Appl.
Mech., 37 83-91 (1971). 20
T. C. Togami, W. E. Baker, and M. J. Forrestal, “A Split Hopkinson Bar
Technique to Evaluate the Performance of Accelerometers,” J. Appl. Mech., 63
353-356 (1996).21
W. Chen, B. Zhang, and M. J. Forrestal, “A Split Hopkinson Bar Technique
for Low-Impedance Materials,” Exp. Mech., 39 81-85 (1999).22
R. J. Christensen, S. R. Swanson, and W. S. Brown,”Split-Hopkinson-Bar
Tests on Rock Under Confining Pressure,” Exp. Mech., 29 508-513 (1972).23
D. J. Frew, M. J. Forrestal, and W. Chen, “Pulse Shaping Techniques for
Testing Brittle Materials with a Split Hopkinson Pressure Bar,” Exp. Mech.,
accepted (2001). 24
M. J. Forrestal, D. J ,Frew, and W. Chen, “The Effect of Sabot Mass on the
Striker Bar for Split Hopkinson Pressure Bar Experiments,” Exp. Mech.,
submitted (2001). 25
W. Chen, Q. Wu, J. H. Kang, and N. A. Winfree, "Compressive Superelastic
Behavior of a NiTi Shape Memory Alloy at Strain Rates of 0.001 to 750 s-1
,"
International Journal of Solids and Structures, in printing (2001). 26
W. Chen, F. Lu, and M. Cheng, “Tension and Compression Tests of Two
Polymers Under Quasi-static and Dynamic Loading” Polymer Testing, in printing
(2001). 27
W. Chen, F. Lu and B. Zhou, “A quartz crystal imbedded split Hopkinson bar
for soft materials,” Experimental Mechanics, 40, (1) pp. 1-6 (2000).
224 Ceramic Armor Materials by Design
USING BAR IMPACT TO DETERMINE DYNAMIC PROPERTIES OF
CERAMICS
Dr. Stephan J. Bless
Institute of Advanced Technology
The University of Texas at Austin
3925 West Braker Lane, Suite 400
Austin, TX 78759
ABSTRACT
Impact onto bars provides a useful means to study the properties of brittle
materials. Measurements may be stress (using piezoresistive gauges) or free
surface velocity. The amplitude of the stress that propagates in the bar is the
largest compressive stress that can be supported by the target material (in a one-
dimensional stress state); however, premature failure due to tension in the impact
zone may limit the peak stress. High-speed photography can be used to reveal the
morphology of the failure – which takes place by longitudinal cracks, transverse
cracks, or failure waves.
INTRODUCTION
It is difficult to develop empirical criteria for compression failure of ceramics
under impact loading. Strength depends on stress rate and stress state. Plate
impact experiments determine the compressive strength at very high rates and full
lateral confinement. Under these conditions, many brittle materials exhibit
considerable ductility.
In fact, high velocity impact induces compressive brittle failure, generally
accompanied by substantial lateral strain. For example, cavity expansion models
of ceramic penetration show that a critical stage of compressive failure takes place
under uniaxial stress [1].
In a plate-impact experiment, the failure stress is usually identified as the
Hugoniot elastic limit HEL. In a bar-impact experiment, the failure stress is
nominally equal to Y, the strength in one-dimensional stress. As discussed
already in this symposium [2], different micromechanical models predict different
relationships between HEL and Y. In some instances, strain-based failure criteria
have also been proposed.
Ceramic Armor Materials by Design 225
To the extent authorized under the laws of the United States of America, all copyright interests in this publication are the propertyof The American Ceramic Society. Any duplication, reproduction, or republication of this publication or any part thereof, withoutthe express written consent of The American Ceramic Society or fee paid to the Copyright Clearance Center, is prohibited.
BAR IMPACT TESTS
The bar impact geometry is shown in Fig. 1. A striker plate or bar impacts a
target bar. The impactor can be the same material as the target, or a hard steel;
more consistent results are usually obtained with metal strikers. If the bar does not
fail, then a stress wave is produced whose amplitude is determined by the
conventional impedance match solution. But the test is usually designed so that
the target bar fails in the vicinity of the impact plane. As in a plate impact test, an
elastic wave then propagates along the bar, and the amplitude of this wave is
equal to the largest stress in the bar before it failed. By definition this is the bar
impact strength, YB. In metals, YB=Y, but in brittle materials YB Y because
failure can initiate from transient impact-induced tensile stresses [3].
Most experiments have been done with round bars. However, tests have also
been performed with square, rectangular, or octagonal bars cut from plates. There
does not appear to be any systematic difference in strength associated with cross
section shape.
The distance from the impact plane to the measurement plane should be about
10 diameters. Shorter distances are probably possible in some materials, but in
[4] it appears that six diameters was too short. Using a layered striker to induce a
ramp wave loading of the target bar seems to enable use of shorter target bars [5].
11
00
.25
98
a LEXAN
GAUGE PLANETARGET ROD
STEEL
LEXAN LEXAN
CERAMIC
OR
Figure 1. Bar impact geometry. The ceramic target rod may be struck by a steel
plate or a ceramic rod.
The stress wave in the bar can be recorded by using an interferometer to
measure the motion at the free end, or by using an embedded stress gauge (see
record in Figure 2). The free surface technique provides the most faithful record
of the rise to peak stress. However, spall failure can occur near the free surface
soon after wave reflection, so if the goal of the test includes measuring the
crushing behavior of the failed column of material, this technique is not useful.
226 Ceramic Armor Materials by Design
0
0
5
5
10
10
15
15
20
20
t [µs]
1 1
2 2
3 3
4 4
5 5
6 6
σ z[G
Pa]
4340
Steel
Gauge failed
AlON 4340 steel
10 mm sq. c.s.101 mm 42 mm
10 mm
51mm
Camera trigger
Figure 2. Gauge records from bar impacts on AlON.
Use of stress gauges in bars has been analyzed and validated by [6]. The best
records are obtained when the brittle bar is backed by a metallic witness bar. The
strength of the witness bar should be greater than the strength of the brittle
material; otherwise, the gauge measures the witness bar strength. For strong
ceramics, it is often necessary to place the gauge between two ceramic bars. This
emplacement, unfortunately, usually results in gauge failure just after the peak
stress signal. The reasons for premature gauge failure are not yet clear.
Bar impact tests have also been performed with lateral confinement. Steel and
tantalum have been used [5]. The measured bar impact strength increases when
the bars are confined. However, the interpretation of the strength is not
straightforward, since the confining stress is rather difficult to determine.
The split Hopkinson bar (SHB) also can be used to measure unconfined
compressive strength. While historically SHB test results have exhibited
considerable scatter, recent progress, reported in this symposium, on wave
shaping holds great promise for achieving reliable strength measurements [7].
Nevertheless, advantages of the bar impact over the Hopkinson bar include less
sample machining, avoiding shape changes, bigger samples, observation of failure
morphology, ability to test very hard materials, ability to study failure
propagation, and easier characterization of the post-failure behavior.
Another attribute of the bar impact test is that there is relatively little scatter.
In several test series, the shot to shot variation in YB has been < 10%. This much
better than usually observed in static unconfined compression tests. Thus, the bar
impact may be an economic means of quality control for ceramics.
Lastly, both the flyer plate and the witness bar can exhibit several different
post-impact appearances. They may be undamaged, indented, or cratered. These
behaviors are indicative of radically different failure modes among various brittle
materials.
Ceramic Armor Materials by Design 227
RESULTS FROM BAR IMPACT TESTS
One of the most useful features of the bar impact test is visual access. High-
speed photos of the dynamic failure process are possible. Several categories of
failures have been observed. The most dramatic observations are of self-
propagating failure waves. Figure 3 shows an example in glass. In such a failure
wave, material apparently is transformed from an intact to a comminuted state.
Transparent material becomes opaque, and there is moderate radial expansion.
Figure 3. Failure wave in silica glass [8].
Failure waves can also be driven by the projectile, in which case standoff
between the impact face and the propagation front is almost contort. Figure 4
shows an example.
Figure 4. Example of a driven failure front, seen in granite [9].
228 Ceramic Armor Materials by Design
The main use of the bar impact tests has been to measure compressive
strength. Table 1 provides a measurement of strength and observations of failure
modes in bar impacts. Less reliable strength data are shown in parentheses.
Table I. Measurements of bar impact strength of ceramics
Material References YB (GPa) Failure
Morphologies
(See Table 2)
Sintered alumina 3, 4, 5, 6, 10,
11, 12
3.6, 4.2 1, 2
Hot pressed alumina 4 4.1 1, 2
TiB2 4 4.9 3
Silicon carbide 4 4.8 3
Boron carbide 4 n/a 1
Soda lime glass 13 2.0 3, 5
Borosilicate glass 2, 4, 11 (1.5) 2.5 3, 4
Aluminum oxynitride 14 4.0 1, 4
Homalite 15 n/a 5
Table II. Failure modes
1. Axial splitting near impact face
2. Transverse faulting away from the impact zone
3. Self-propagating failure wave
4. Failure along central axis
5. Driven failure wave
In addition to these studies of compressive failure, impact-induced tensile
failure has been studied in ceramic bars by [16].
EMERGING NEW CAPABILITIES
At The University of Texas, we are expanding the repertory of bar impact test
techniques with experiments on a surrogate material – homalite. Homalite is a
brittle thermoset plastic, 1.23 g/cm3, Y = 0.15 GPa.
One promising new area is variations in bar cross section. By using tapered
bars, we achieve a condition in which the stress increases continuously along the
bar. This may avoid the problem of premature failure due to impact transients.
Figure 5 is a photograph of a test using a tapered homalite bar.
Ceramic Armor Materials by Design 229
Figure 5. Photograph of impact induced failure in a tapered homalite bar.
Recovery of fractured material is also an area that has received little attention
except [11], but is potentially very useful. We have recently developed a
technique in which a 10-mm bar is sleeved in plastic. The particles are recovered
within the sleeve, and they can be correlated with fracture zones seen in
photographs of the bars. These include the comminuted region associated with
failure fronts, faulted material produced relatively late in the impact, and spall
fragments formed at the rear surface. So far in homalite, we have not seen
evidence of fractal behavior, nor of ductile flow. Rather, we observed that in the
comminuted region there is a well defined smallest particle size, about 20
microns. Particles from elsewhere along the bar are much larger, ranging in size
up to several mm.
It has also become clear that there are phenomena associated with bar impact
that are not yet understood. Figure 6, for example, shows precursor cracks that
will lead rapidly to an isolated zone of comminuted material. This type of failure,
which is not due to simple inversion of the loading pulse at the free end, has not
yet been modeled.
Figure 6. Spall failures developing in the interior of a homalite bar [13].
230 Ceramic Armor Materials by Design
ACKNOWLEDGEMENTS
The research reported in this document was performed in connection with
Contract number DAAD17-01-D-0001 with the U.S. Army Research Laboratory.
The views and conclusions contained in this document are those of the authors
and should not be interpreted as presenting the official policies or position, either
expressed or implied, of the U.S. Army Research Laboratory or the U.S.
Government unless so designated by other authorized documents. Citation of
manufacturer’s or trade names does not constitute an official endorsement or
approval of the use thereof. The U.S. Government is authorized to reproduce and
distribute reprints for Government purposes notwithstanding any copyright
notation hereon. Additional data were provided by Rod Russell at The University
of Texas.
REFERENCES
[1] S. Satapathy and S.J. Bless, “Cavity Expansion Resistance of Brittle
Materials Obeying a Two Curve Pressure Shear Behavior,” Journal of Applied
Physics, 88 [1] 4004-4012 (1999).
[2] G.I. Kanel and S.J. Bless, “Compressive Fracture of Brittle Solids under
Shock-Wave Loading,” Int’l Conference on Advanced Ceramics and Glasses
(PacRim IV), Nov. 4-8, 2001, to be published by American Ceramics Society
(2002).
[3] C.H.M. Simha, S.J. Bless, and A. Bedford, “What is the Peak Stress in the
Ceramic Bar Impact Experiment?”; pp.615-618 in Shock Compression of
Condensed Matter – 1999. Edited by M.D. Furnish, L.C. Chhabildas, and R.S.
Hixson. American Institute of Physics, 2000.
[4] N.S. Brar and S.J. Bless, “Dynamic Fracture and Failure Mechanisms of
Ceramic Bars,” On Shock-Wave and High-Strain-Rate Phenomena in Materials
(EXPLOMET 90), Aug. 12-17, 1990, to be published.
[5] L.C. Chhabildas, M.D. Furnish, and D.E. Grady, “Impact of Alumina Rods
– A Computational and Experimental Study,” J. Phys. IV, 4 [C3] 137-143 (1997).
[6] Z. Rosenberg, P. Partom, and B. Keren, “Gauge Factor of Manganin under
Axial Stress Conditions,” J. Appl. Phys., 54 2824-2825 (1983).
[7] W. Chen, B. Song, D.J. Frew, and M.J. Forrestal, “Recent Developments
in Split Hopkinson Pressure Bar Testing,” Int’l Conf. On Advanced Ceramics and
Glasses (PacRim IV), Nov. 4-8, 2001, to be published by American Ceramics
Society (2002).
[8] S.J. Bless, N.S. Brar, G. Kanel, and Z. Rosenberg, “Failure Waves in
Glass,” J. Am. Ceram. Society, 75 [1] 1002-1004 (1992).
[9] L. Glenn and W. Janach, “Failure of Granite Cylinders under Impact
Loading,” Int’l. J. Fracture, 13 [1] 301-317 (1977).
Ceramic Armor Materials by Design 231
[10] J.U. Cazamias, B. Reinhart, C. Konrad, L.C. Chhabildas, and S.J. Bless,
“Bar Impact Tests on Alumina (AD995),” Shock Compression of Condensed
Matter 2001, to be published by American Institute of Physics, 2002.
[11] H.D. Espinosa, Y. Xu, and N.S. Brar, “Micromechanics of Failure Waves
in Glass: I, Experiments,” J. Am. Ceram. Soc. 80 [1] 2061-73 (1997).
[12] J.L. Wise and D.E. Grady, “Dynamic, Multi-axial Impact Response of
Confined and Unconfined Ceramic Rods”; pp. 777-780 in High-Pressure Science
and Technology. Edited by S.C. Schmidt et al. 1993. AIP Conference Proceedings
309 (1994).
[13]. N.H. Murray, N.K. Bourne, J.E. Field, and Z. Rosenberg, “Symmetrical
Taylor Impact of Glass Bars,” Shock Compression of Condensed Matter – 1997,
American Institute of Physics (1998).
[14] J.U. Cazamias, P.S. Fiske, and S.J. Bless, “The Hugoniot Elastic Limit of
AlON,” Shock Compression of Condensed Matter-2001, to be published by
American Institute of Physics (2002).
[15] R. Russell, S. Bless, and T. Beno, “Impact Induced Failure
Phenomenology in Homalite Bars,” Shock Compression of Condensed Matter-
2001, to be published by American Institute of Physics (2002).
[16] J. Najar, “Dynamic Tensile Fracture Phenomena at Wave Propagation in
Ceramic Bars,” J. Physics IV 1 [C8] 647-652 (1994).
232 Ceramic Armor Materials by Design
SHOCK COMPRESSION AND RELEASE PROPERTIES OF COORS AD995
ALUMINA
William D. Reinhart Lalit C. Chhabildas
Sandia National Laboratories Sandia National Laboratories
Weapons Science Applications Weapons Science Applications
PO Box 5800 PO box 5800
New Mexico, 87185-1181 New Mexico, 87185-1181
Dennis E. Grady Tsutomu Mashimo
Applied Research Associates, Inc. High Energy Rate Laboratory
4300 San Mateo Blvd. NE. Kumamoto University
Albuquerque, New Mexico, 87110 Kumamoto 860, Japan
ABSTRACT
An investigation of the shock compression, recompression and decompression
properties of Coors AD995 alumina (aluminum oxide) ceramic and single crystal
sapphire has been conducted. Well-controlled, planar impact experiments have
been performed in which stationary targets are impacted by ceramic plates to
pressures exceeding 100 GPa. In this study of Coors AD995 ceramic and single
crystal sapphire, dynamic material property data is obtained utilizing gun loading
techniques and high-resolution velocity interferometric tools. Techniques used to
determine the dynamic compression, recompression, and release behavior are
summarized herein.
INTRODUCTION
Ceramics in general have repeatedly demonstrated to be an effective armor
material due to its high dynamic yield strength compared to metals. However, like
many brittle materials ceramics are weak in tension as evidenced by the low spall
strength of the materials. There is also evidence that the transient strength of
many of these materials degrades - which has led to the concept of the existence
of failure waves in materials and is indicative of damaged material even when
under compression. Kanel et al [1] was the first to show conclusively from his
experiments on glass that the dynamic yield strength of glass decreases as the
shock dwell time increased. There is also evidence of the dynamic yield strength
Ceramic Armor Materials by Design 233
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degradation in boron carbide ceramic [2] when shocked above its Hugoniot elastic
limit.
Shock experiments are confined uniaxial strain experiments and are generally
referred to as the Hugoniot state of the material. To estimate the dynamic yield
strength of the material, the Hugoniot state is compared to a hydrostat - which is
determined by extrapolating the stress-strain behavior of the material determined
at lower hydrostatic pressures. Based on Von-Mises yield criteria the difference
between the Hugoniot stress and the hydrostatic pressure curve is defined as two-
thirds the dynamic yield strength. If this difference is (1) independent of the
shock-loading stress then the material exhibits elastic-perfectly plastic behavior,
(2) changing with increasing stress then the material exhibits a pressure-
dependent yield strength. An increase in yield strength may be attributed to many
factors such as rate-dependence and or a pressure dependent yield behavior, while
a decrease would be related to a softening behavior resulting from heterogeneous
deformation process and or from damage resulting from shock compression.
One of the objectives of the present study is to investigate the possibility of
determining, dynamically, the shock-hydrostat for ceramics. This technique has
been previously applied to investigate metals [3,4], and in particular has been
used extensively to characterize the strength properties of 6061-T6 aluminum and
tungsten in the shocked state. The method employs re-shock and release
experiments to be conducted from the same Hugoniot stress state to
experimentally evaluate the departure of the initial loading stress state from an
elastic plastic behavior. The asymmetry in the reloading and release path is then
used to determine the shock-hydrostat. In this study, well-controlled impact
experiments are performed on smooth-bore guns, and velocity interferometric
diagnostics [5] are used to acquire high-resolution shock compression, and
subsequent recompression or release data on alumina.
Aluminum oxide (Al2O3) is a widely used commercial ceramic because of its
useful electrical, and mechanical properties. It also has good optical properties
when used as a single crystal, commonly known as sapphire. In shock
experiments single crystal sapphire has been used as laser-interferometer
windows [6]. Extensive shock-Hugoniot equation-of-state studies have been
performed on aluminum-oxide primarily because of its wide applicability as an
armor ceramic. Sapphire was included in this study because it is the single crystal
form of Al2O3 and is the building block at a granular level. In this study, only
reshock and release experiments at 27 and 44 GPa on alumina are highlighted,
however, the Hugoniot experiments are reported to a peak stress level of over 1
Mbar (for alumina and sapphire).
234 Ceramic Armor Materials by Design
MATERIAL
The aluminum oxide (Al2O3) used in this study is referred to as Coors AD995.
Its composition consists of 99.5% alumina and the remainder of the material is
aluminosilicate glass. The density of the material (Al2O3) was 3.89 g/cm3 and the
average longitudinal and shear wave speed was 10.56 km/s and 6.24 km/s
respectively. This yields an estimate of 7.71 km/s, 0.234 and 231.7 GPa for the
bulk wave velocity, Poisson’s ratio and the bulk modulus, respectively. Sapphire,
which is the single crystal form of Al2O3, has a rhombohedral-hexagonal crystal
structure with close-packed oxygen ions. Both c-axis and a-axis crystals were
used in this study. The elastic longitudinal wave speed for the c-axis and a-axis
crystal was determined to be 11.19 km/s [6] and 11.17 km/s [7]. The density of
the sapphire crystals used in this study was 3.98 gm/cm3.
EXPERIMENTAL METHOD
Compressive shock, reshock and release waves are produced in aluminum
oxide and sapphire with a single stage powder gun and a two-stage light gas gun.
The experimental configuration used for this study for both the powder gun and
the two-stage light gas gun is shown in Figure 1. The powder gun has an 89 mm
bore diameter and achieves impact velocities exceeding 2.3 km/s, while the two-
stage light gas gun utilizes a 28 mm bore diameter with projectile velocities
approaching 8 km/s. The powder gun projectile velocity is measured by three
electrical self-shorting pins, which are mounted on the target fixture, to accuracy
better than 0.5%. Additional electrical pins are incorporated to measure impact
planarity (typically about 1-2 milli-radians) and provide triggers to diagnostic
equipment. The two stage light gas gun incorporates a projectile velocity
measuring system call the Optical Beam Reflector [8] (OBR) that accurately
Ceramic
Low or High
Impedance Backing
Lithium
Fluoride
Window
Electrical Self
Shorting Pins
Projectile
VISAR
Ceramic
Low or High
Impedance Backing
Lithium
Fluoride
Window
Electrical Self
Shorting Pins
Projectile
VISAR
Figure 1. Experimental Configuration
Ceramic Armor Materials by Design 235
measures projectile velocity to better than 0.2%. As with the target on the powder
gun, electrical pins are also used for impact planarity and diagnostic triggering.
SPT-2:
SPT-1:
SPT-4:
0.0
0.5
1.0
1.5
2.0
2.5
3.0
3.5
Win
do
w V
elo
city
(k
m/s
)
0.25 0.30 0.35 0.40
Arbitrary time ( s)
CE58 ALRL3
ALRS2
0.5 1.0 1.5 2.0 2.5
Arbitrary time ( s)
0.0
0.5
1.0
1.5
2.0
Win
do
w V
elo
city
(k
m/s
)
SPT-2:
SPT-1:
SPT-4:
0.0
0.5
1.0
1.5
2.0
2.5
3.0
3.5
Win
do
w V
elo
city
(k
m/s
)
0.25 0.30 0.35 0.40
Arbitrary time ( s)
CE58 ALRL3
ALRS2
0.5 1.0 1.5 2.0 2.5
Arbitrary time ( s)
0.0
0.5
1.0
1.5
2.0
Win
do
w V
elo
city
(k
m/s
)
ALRS1ALRS1a)
.
b).
SPT-3:CE60
SPT-3:CE60
Figure 2. (a) Window/interface particle velocity profiles for Al203, and (b) c-
axis and a-axis sapphire, respectively.
In Figure 1, the projectile is faced with the ceramic Coors AD995 or sapphire
and is backed with either a foam disk of low shock impedance, or a high shock
impedance material, tantalum, for reshock experiments. The target configuration
in Figure 1, will have a alumina (or sapphire) ceramic disk similar to that
mounted on the projectile and a single crystal lithium-fluoride is bonded with
epoxy to the back of the ceramic sample. The lithium-fluoride is an optical
quality disk, lapped and polished and is typically flat to within a few bands of
sodium light. One surface of the lithium-fluoride is diffused and approximately
100nm of aluminum is vapor deposited on the lapped surface before being glued
to the alumina disk. The particle velocity histories resulting from impact were
measured at the target/lithium-fluoride window [9] interface using a velocity
interferometer, VISAR [5]. The Doppler shifted interference fringes measured
with the VISAR are converted to a time-resolved velocity history and are shown
in Figures 2a and 2b for the experiments on alumina and sapphire, respectively.
The amplitude resolution is approximately
2% per fringe and typically two to three
fringes are achieved in the interface
acceleration resulting from the compressive
shock front.
EXPERIMENTAL RESULTS
Elastic Waves
The impact conditions for the
experiments in the current study are
summarized in Table I. The wave profiles
shown in Figures 2 are used to determine the
Elastic Shock
Elastic-Plastic
Deform ation Hugoniot State
M id-Point of shock
Elastic Shock
Elastic-Plastic
Deform ation Hugoniot State
M id-Point of shock
Figure 3. Shock wave profile
in alumina traversing from right
to left.
236 Ceramic Armor Materials by Design
Hugoniot properties of the ceramic. Hugoniot refers to the peak stress states
achieved in the shock compression process. Figure 3 illustrates the particle
velocity (stress) wave profile traversing toward the left through the ceramic. In
this illustration, material to the left of the elastic shock front is undisturbed. The
leading edge of the precursor wave is used as a fiducial for the analysis in this
study. For alumina, the leading edge of the elastic wave traverses at the elastic
longitudinal wave speed of 10.74 km/s, a value that has been determined on
earlier studies in the alumina [10]. For sapphire the leading edge of the wave
traverses at 11.7 km/s, and is consistent with the linear elastic behavior reported
in the literature [6]. The Hugoniot elastic limit stress, ( hel), is determined using
the relation:
hel = ( o Cl ue ), (1)
where o is the initial density of the ceramic, Cl the elastic longitudinal wave
speed, and ue is the in-material particle velocity measurement prior to transition
to a plastic wave.
Table I. Summary of Impact Conditions
Exp.
No.
Impact
Velocity
(km/s)
Target
Thickness
(mm)
Impactor
Thickness
(mm)
ue
(km/s) e HEL
(GPa) e
Elastic
Strain
CE57 1.019 10.006 5.019 0.153 6.41 0.0143
CE58 1.572 10.008 5.008 0.154 6.44 0.0143
CE59 2.030 10.007 5.013 0.150 6.26 0.0140
CE60 2.329 9.998 5.005 0.178 7.44 0.0166
CE61 0.561 9.998 5.013 0.148 6.18 0.0136
CE62 2.211 9.987 5.005 0.175 7.32 0.0163
CE63 2.062 9.987 4.989 0.162 6.75 0.0151
ALRL3 2.158 7.988 3.070 0.169 7.06 0.0157
ALRS11
2.185 6.335 0.4992 3 3 3
ALRS2 2.208 6.337 4.211 0.162 6.78 0.0151
SAPT14
4.220 3.195 12.697 0.506 24.17 0.0432
SAPT24
4.431 3.193 12.708 0.527 24.48 0.0451
SAPT35
3.28 3.387 1.4665
0.5 24.00 0.0427
SAPT45
3.30 3.422 1.4875
0.532 25.50 0.04551 Reverse ballistic experiment 2 Aluminum Buffer for Window experiment
Table I. Summary of Impact Conditions
Exp.
No.
Impact
Velocity
(km/s)
Target
Thickness
(mm)
Impactor
Thickness
(mm)
ue
(km/s) e HEL
(GPa) e
Elastic
Strain
CE57 1.019 10.006 5.019 0.153 6.41 0.0143
CE58 1.572 10.008 5.008 0.154 6.44 0.0143
CE59 2.030 10.007 5.013 0.150 6.26 0.0140
CE60 2.329 9.998 5.005 0.178 7.44 0.0166
CE61 0.561 9.998 5.013 0.148 6.18 0.0136
CE62 2.211 9.987 5.005 0.175 7.32 0.0163
CE63 2.062 9.987 4.989 0.162 6.75 0.0151
ALRL3 2.158 7.988 3.070 0.169 7.06 0.0157
ALRS11
2.185 6.335 0.4992 3 3 3
ALRS2 2.208 6.337 4.211 0.162 6.78 0.0151
SAPT14
4.220 3.195 12.697 0.506 24.17 0.0432
SAPT24
4.431 3.193 12.708 0.527 24.48 0.0451
SAPT35
3.28 3.387 1.4665
0.5 24.00 0.0427
SAPT45
3.30 3.422 1.4875
0.532 25.50 0.04551 Reverse ballistic experiment 2 Aluminum Buffer for Window experiment 3 Elastic values not determined 4 C-cut single crystal sapphire 5 Tungsten ( = 19.2 g/cm3) used to Impact A-axis sapphire crystal.
Ceramic Armor Materials by Design 237
Figure 4. Hugoniot results. a)
shock–velocity vs material
velocity b) stress vs. material
velocity c) stress vs strain
0
20
40
60
80
100
120
0.00 0.05 0.10 0.15 0.20 0.2Strain ( )
Str
ess (
GP
a)
C - cut Sapphire
Aluminum Oxide
A - cut Sapphire
0
20
40
60
80
100
120
0.0 0.5 1.0 1.5 2.0 2.5Particle Velocity (km/s)
Str
ess (
GP
a)
C- cut Sapphire
Aluminum Oxide
A - cut Sapphire
6
8
10
12
14
0.0 0.5 1.0 1.5 2.0 2.5
Particle Velocity (km/s)
Sh
ock V
elo
cit
y (
km
/s)
C - cut Sapphire
Aluminum Oxide
HEL-Sapphire
HEL-Al203
A - cut Sapphire
The in-material particle velocity is
determined through the impedance
matching relation:
ue = uw (Zw + Zm) / 2 Zm (2)
where uw, Zw, and Zm are the measured
velocity in the window material, and the
shock impedance of the lithium-fluoride
window and the ceramic material
respectively. The shock impedance of
the material is defined as the product of
its density and the shock velocity. The
equation of state of lithium-fluoride [10]
is used to calculate its shock impedance,
while the elastic shock impedance of the
alumina or sapphire is the product of o
and Cl, their respective density and the
elastic wave speed. The elastic strain, e,
is calculated using ue/Cl . The stress,
strain and the particle velocity results at
the elastic limit for the series of
experiments reported in this paper are
tabulated in Table I and shown in Figure
4.
Plastic Waves
The planar impact produces a
compressive wave of uniaxial strain,
which propagates across the target
specimen and into the lithium-fluoride
window. The measured velocity exhibits
a two-wave structure. The subsequent
structure following the elastic precursor
represents pressure hardening of the
material and this two-wave structure is
the result of a transition from elastic to
plastic deformation. As compression
within the shock increases during the
shock loading process, shear stresses
Figure 4. Hugoniot results. a)
shock–velocity vs material
velocity b) stress vs. material
velocity c) stress vs strain
0
20
40
60
80
100
120
0.00 0.05 0.10 0.15 0.20 0.25Strain ( )
Str
ess (
GP
a)
C - cut Sapphire
Aluminum Oxide
A - cut Sapphire
0
20
40
60
80
100
120
0.0 0.5 1.0 1.5 2.0 2.5Particle Velocity (km/s)
Str
ess (
GP
a)
C- cut Sapphire
Aluminum Oxide
A - cut Sapphire
6
8
10
12
14
0.0 0.5 1.0 1.5 2.0 2.5
Particle Velocity (km/s)
Sh
ock V
elo
cit
y (
km
/s)
C - cut Sapphire
Aluminum Oxide
HEL-Sapphire
HEL-Al203
A - cut Sapphire
238 Ceramic Armor Materials by Design
will exceed the critical strength of the material (HEL) and plastic deformation
occurs in the observed second wave.
Because finite rise times are measured for the plastic wave, the plastic-
wave velocity, Usp, is taken at the center of the wave (Figure 3) and the
corresponding wave speed is given in Table II. Where symmetric impact
techniques are used, the particle (material) velocity, uph, behind the shock front is
exactly one-half the impact velocity. The Hugoniot stress, ph, and strain, ph,
behind the plastic-wave front are estimated using the following relations:
ph = e + [ o Usp (uph – ue)] (3)
ph = e + (uph – ue) / Usp (4)
The summary of all Hugoniot data in Tables I and II, are shown plotted in Figure
4, represents the study on Coors AD995 and agrees well with previous work
[10,12,13,14] of this material. It should be noted that the HEL of this material is
about 6.0 to 7.5 GPa while the HEL of the single crystal material is about 20 to
24 GPa. As indicated in Table II, the current studies span over the stress regime
of 18 to 100 GPa.
Table II. Hugoniot Summary
Exp. No. Usp
(km/s)
uph
(km/s) ph
(GPa) ph
Total Strain
CE57 8.38 0.510 18.02 0.0568
CE58 8.40 0.787 27.11 0.0896
CE59 8.49 1.015 34.83 0.1159
CE60 8.81 1.165 42.30 0.1286
CE61 9.18 0.281 11.01 0.0280
CE62 9.48 1.452 54.69 0.1502
CE63 9.72 1.457 55.83 0.1477
ALRL3 8.94 1.079 38.70 0.1176
ALRS1 8.86 0.762 26.27 0.0860
ALRS2 8.93 1.104 40.38 0.1178
SAPT1 10.54 2.110 90.62 0.1953
SAPT2 10.72 2.215 96.31 0.2025
SAPT3 10.92 2.285 101.3 0.2062
SAPT4 10.94 2.390 106.2 0.2153
Ceramic Armor Materials by Design 239
Off-Hugoniot States
Most of the experiments conducted in this investigation, evaluate the off-
Hugoniot states of Coors AD995 alumina as it is allowed to release from
compression. Experiments ALRS1 and ALRS2 are the experiments on Coors
AD995 alumina as it is further recompressed from its original shocked Hugoniot
state. An incremental form of the conservation equations given by the relations:
= o c u (5)
= u/c (6)
is used to estimate the final released or reshocked stress, and strain, ,
respectively. The Lagrangian-wave velocity, c, corresponds to the material
particle velocity change, u. This is indicated in the x-t diagram in Figure 5a.
Although the figure emphasizes the reshock experimental configuration, the same
concept is true for a release diagram. The backing to the ceramic impactor is
replaced by carbon foam, a low shock impedance material, so that a reflected
release wave will propagate in the impactor material when the shock arrives at the
Ta
nta
lum
Ceramic
Win
dow
Shock
Elasti
c-Res
hock
Plasti
c-Res
hock
Time
X
Ta
nta
lum
Ceramic
Win
dow
Shock
Elasti
c-Res
hock
Plasti
c-Res
hock
Time
X
Tan
talu
m
Ceramic Ceramic
Win
dow
ElasticPlasti
c
Release Waves
Elastic-
Reshock
Plasti
c-Res
hock
te
tp
tf
tp
Different Wave Velocities
resulting from wave
interactions
High Stress:
SampleWindow Stress
Us effec
tive
Time
X
Tan
talu
m
Ceramic Ceramic
Win
dow
ElasticPlasti
c
Release Waves
Elastic-
Reshock
Plasti
c-Res
hock
te
tp
tf
tp
Different Wave Velocities
resulting from wave
interactions
High Stress:
SampleWindow Stress
Us effec
tive
Time
X
b)a)
Figure 5. X-t diagram typical experiments. a) depiction of symmetric impact
with wave interactions, b). alternate technique where by wave interactions are
eliminated.
Ta
nta
lum
Ceramic
Win
dow
Shock
Elasti
c-Res
hock
Plasti
c-Res
hock
Time
X
Ta
nta
lum
Ceramic
Win
dow
Shock
Elasti
c-Res
hock
Plasti
c-Res
hock
Time
X
Tan
talu
m
Ceramic Ceramic
Win
dow
ElasticPlasti
c
Release Waves
Elastic-
Reshock
Plasti
c-Res
hock
te
tp
tf
tp
Different Wave Velocities
resulting from wave
interactions
High Stress:
SampleWindow Stress
Us effec
tive
Time
X
Tan
talu
m
Ceramic Ceramic
Win
dow
ElasticPlasti
c
Release Waves
Elastic-
Reshock
Plasti
c-Res
hock
te
tp
tf
tp
Different Wave Velocities
resulting from wave
interactions
High Stress:
SampleWindow Stress
Us effec
tive
Time
X
b)a)
240 Ceramic Armor Materials by Design
0
10
20
30
40
50
60
0.00 0.04 0.08 0.12 0.16
Strain ( )
Loading Path
Hugoniot Stress
Reshock
elastic
0
10
20
30
40
50
60
0.00 0.04 0.08 0.12 0.16
Strain ( )
Loading Path
Hugoniot Stress
Reshock
elastic
Str
es
s (
GP
a)
Str
es
s (
GP
a)
The results of companion
reshock and release experiments
conducted at approximately
27 GPa and 40 GPa are shown in
Figures 6 and 7, respectively.
The results of these experiments
will be highlighted in this paper.
Also shown in the figure is a
calculated hydrostat for the
alumina based on Murnaghan
equation of state where the bulk
modulus of Coors AD995
alumina is used. The bulk
modulus is based on the
ultrasonic sound speed
measurements on the samples. A value of four (4) is used for the pressure
backing/sample interface.
The analysis is
approximated by
representing the finite rise-
time of the shock in the
impactor as a single shock
wave traversing at an
effective shock velocity,
calculated using the
relation, Ueff = ph/( ouph),
where the stress ph and
particle velocity uph are the
first shocked states. The
Lagrangian-wave velocity
is estimated from the time difference between the arrival of the leading edge of
the release/reshock wave at the window/sample interface and the time at which
the effective shock arrives at the back surface of the impactor (Finite rise times
upon reshock are measured because an elastic-plastic wave is observed). As
indicated in figure 5, the release fan or the reshock is also perturbed by the
reflected release wave that emanates at the target window interface. Making the
ratio of the target sample dimension to the impactor dimension large, confines the
interaction zone towards the window interface, and minimizes this perturbation.
An alternate technique (Figure 5b) is to impact the window directly so that all
wave interactions are eliminated. This is specifically done in the reshock
experiment ALRS1.
0
10
20
30
40S
tress (
GP
a)
0.00 0.02 0.04 0.06 0.08 0.10 0.12
Strain ( )
Loading Path
Hugoniot Stress
Calculated Hydrosta
t
UnL
oadin
gPath
s
Reshock
0
10
20
30
40S
tress (
GP
a)
0.00 0.02 0.04 0.06 0.08 0.10 0.12
Strain ( )
Loading Path
Hugoniot Stress
Calculated Hydrosta
t
UnL
oadin
gPath
s
Reshock
elastic
Figure 6. Stress-strain plots for ALRS1 and
CE58 depicting loading and unloading paths.
UnL
oading P
ath
UnL
oading P
ath
Figure 7. Stress-Strain plots for
ALRS2 and CE60 depicting
compression, recompression and
unloading behavior.
Ceramic Armor Materials by Design 241
derivative of the bulk modulus and is based on high-pressure x-ray diffraction
data on single-crystal sapphire [15].
DISCUSSIONS
In this study, we are reporting the shock Hugoniot of Coors AD995 alumina to
approximately 60 GPa, and that of sapphire to stresses over 100 GPa. The
alumina used in this study is the same batch of material used in an earlier
investigation [10,16]. Results from previous investigations [12,13,14] are not
reported herein mainly because the material studied may not be quite the same –
even though the results do suggest a good agreement with past studies on other
types of alumina.
Elastic- plastic waves
The Hugoniot elastic limit of Coors AD995 reported in Table I vary from
about 6.0 GPa to 7.5 GPa, while the corresponding values for single crystal
sapphire is approximately 24 GPa. This is consistent with previous studies on
single crystals in other materials [16,17] that report higher elastic limits than
polycrystalline materials. A good example is the previous study in crystalline
quartz [18] that has indicated different values of HEL for different crystal
orientation. This is because different slip systems are activated during the
dynamic yielding process. In a polycrystalline material, all these different slip
systems are randomly distributed and hence the results will be dominated by the
weaker slip systems (such as grain boundaries, presence of glass) that could yield
at a lower stress. The Hugoniot elastic limit in alumina will be overdriven when
the shock velocity measurement exceeds the elastic-wave velocity of 10.74 km/s.
Another interesting feature observed in this study is the transition from
elastic yielding to plastic deformation for the two different materials. The post
yielding process in polycrystalline alumina is considerably ramped suggesting a
work hardening type of process or heterogeneous yielding from different slip
systems followed by a plastic-wave that has a finite rise time even up to stresses
approaching 60 GPa. For single-crystal sapphire, however, the transition to the
plastic deformation state is very different. The post yielding process appears to be
either constant or suggests a decrease prior to the arrival of the plastic wave. This
decrease could either be due to elastic precursor decay or may be indicative of a
softening behavior in the single crystal. Also, the rise-time of the plastic-wave for
single crystal sapphire is extremely rapid. It should also be noted that even at
100 GPa, the elastic limit is not overdriven. In this instance, the elastic limit will
be overdriven when shock velocities exceed 11.6 km/s. The interface particle-
velocity measurement for single-crystal sapphire suggests fluctuations behind the
shock front and is presumably a result of heterogeneous deformation process in
the single crystal.
242 Ceramic Armor Materials by Design
Shock Velocity vs Particle Velocity
The variation in shock velocity vs. particle velocity for both materials is
shown in Figure 4(a). As indicated in the figure the elastic limit for Coors AD995
and single crystal sapphire extend up to particle velocity of ~ 0.15 and 0.5,
respectively. The shock-velocity measurements below a particle-velocity
measurement of 0.5 km/s indicate a quasi-elastic behavior, for polycrystalline
alumina, indicates that the dynamic yielding process may not be totally complete.
The shock-velocity measurements above a particle velocity of 0.5 km/s suggest
that the yielding process may be very nearly complete. As indicated in Figure
4(a), the observed shock-velocity vs. particle-velocity variation for single-crystal
sapphire experiments at Mbar pressures appear to be consistent with the
measurements depicted for alumina at stresses up to 50 GPa. However, for a
conclusive interpretation there should be overlapping experiments in the same
shock-velocity/particle-velocity regime for the two different materials. A least
squares fit to the shock vs. particle velocity data beyond the elastic limit would
yield the relation:
Us (km/s) = 1.675up + 7.027 (7)
Stress vs. Strain
The values for stress vs. strain for both the materials tabulated in Table I
and II are shown plotted in Figure 4c. The HEL for the single crystal sapphire is
larger than the HEL of Coors AD995 alumina. As indicated in the figure, the
volume compression has being determined to stresses above 100 GPa and to about
22% strain. The experiments do suggest that the single-crystal sapphire
compression appears to be consistent with that of polycrystalline Coors AD995
alumina at high stresses. This implies that either the shear stress in the shocked
state are similar for both materials or that both materials have collapsed to the
hydrostat. It should also be noted that in the shocked state, melting cannot be
ruled out at about 100 GPa. It would, however, be desirable to have overlapping
experiments at the same stress. This also provides an incentive to conduct re-
shock and release experiments in the future for both materials at megabar stresses
to verify these assumptions.
Stress and Wave Speed vs. Particle Velocity
The stress vs. particle velocity behavior is shown in Figure 4b. As in
figures 4a and 4b, the figure shows a similar behavior as discussed above, namely
higher HEL’s for the single crystal sapphire when compared to the polycrystalline
material alumina. A least squares fit to the stress (GPa) vs. particle-velocity
(km/s) data beyond the elastic limit would yield the relation:
Ceramic Armor Materials by Design 243
ph = uph uph2
(8)
The Lagrangian wave speed, Cl, in the material can be obtained by using the
relation:
Cl = (1/ o) (d ph / duph) (9)
which yields;
Cl (km/s) = 7.45 + 3.37 uph (10)
Equation (10) is plotted as a function of particle velocity in Figure 8 as a solid
line. Based on a constant Poisson’s ratio of 0.247, a Lagrangian elastic wave-
speed is calculated and is also shown (as a dashed line) in Figure 8. The
experimentally determined stress vs. particle-velocity relation can be used to
determine the Lagrangian bulk and elastic wave speed in the material (to obtain
the Eulerian wave speed one can use the relation Ce = ( o/ Cl.). As indicated in
Figure 5, we can also calculate the speed at which the leading edge of the release
or reshock wave traverses in the shocked material. The Lagrangian wave speeds
6
10
14
18
22
0.5 1.0 1.5 2.0 2.5Particle Velocity (km/s)
Lag
ran
gia
n W
ave S
peed
(km
/s)
Experiments
Experimentally Calculated
Elastic Wave Speed Assuming
Constant Poisson's Ration
Figure 8. Variation of Lagrangian elastic and bulk wave
speed as a function of particle velocity.
244 Ceramic Armor Materials by Design
of the leading edge of the wave obtained from these experiments are shown as
triangles (fitted with a dashed line) in Figure 8. As indicated in the figure, the
wave-speed measurements obtained from the off-Hugoniot experiments indicate
that the material response remains elastic up to a particle velocity of 1.5 km/s. At
a particle velocity of 2.2 km/s it collapses to the hydrostat. This also explains why
the single-crystal sapphire stress-strain compression data at about a Megabar is
consistent with the polycrystalline stress-strain compression results (see figure
4c).
Reshock and Release States
The results of companion reshock and release experiments conducted at
approximately 27 GPa and 40 GPa are shown in Figures 6 and 7, respectively.
Also shown in Figure 6, is a calculated hydrostat for the alumina based on
Murnaghan equation of state where the bulk modulus of Coors AD995 alumina is
used. In both these experiments the leading edge of the reshock or release wave
traverses at an elastic wave speed (see Figure 8). The release path exhibits an
elastic release from the initial shocked state. The reloading path shows precursor
elastic recompression and the final reshocked state to about 37 GPa. This reshock
state lies above the calculated hydrostat and Hugoniot states. Similar results are
obtained for the reshock and release experiment at approximately 40 GPa. In this
experiment, within the experimental uncertainties, extrapolation of the static
hydrostat to very high pressures and the current experiments, the shocked state
appears to be on the hydrostat. During recompression from about 40 to 57 GPa,
the recompression wave exhibits an elastic recompression. The leading edge of
the release-wave traverses at an elastic wave velocity and as evidenced by the
elastic release. This is indicative of the loss of shear strength in the material; this
phenomena has been observed previously by Asay and Chhabildas in 6061-T6
aluminum, [2] and by Kanel [1] in glass. The damaged material in alumina
resulting from shock compression is presumed to cause the shear strength loss.
Therefore, the material is exhibiting strength recovery during the recompression
process. This phenomena has also been observed in Coors AD995 alumina even
at its Hugoniot elastic limit [17]. Although the shocked state of the material
exhibited in Figure 6 lies above the hydrostat – this technique can also be used to
experimentally determine a shock hydrostat [3] at very high dynamic stresses and
will be the subject of future discussions. This technique is anticipated to be more
accurate than extrapolating the hydrostatic data because the dynamic hydrostat
will be the mean value of the reshock and release end states from a common
Hugoniot state.
Ceramic Armor Materials by Design 245
SUMMARY
Shock compression, recompression and decompression properties are
summarized. Well-controlled, planar impact experiments have been performed
on Coors AD995 ceramic and single-crystal sapphire ceramic plates to pressures
exceeding 100 GPa. In this study of Coors AD995 ceramic and single-crystal
sapphire, dynamic material property data are obtained utilizing gun loading
techniques combined with high-resolution velocity interferometric tools.
Substantial experimental data on the dynamic response of alumina and
sapphire exist. Results of these studies are unique in that they are probing the
strength properties of the material in the shocked state using reshock and release
test methodologies. This has allowed wave speed/sound speed measurements to
stresses above 1 Mbar. Strength loss in the Hugoniot state is evidenced by
precursor elastic compression in the recompression process. This should allow
the development of damage models needed for use as material models in
computational codes. This technique also can yield the dynamic shock hydrostat
to pressure in excess of 1 Mbar using experimental methods and will be the
subject for discussions in future investigations.
REFERENCES
1. G. I. Kanel, S.V. Rasorenov, V. E. Fortov, The Failure Waves and
Spallations in Homogeneous Brittle Materials, Shock Compression of Condensed
Matter, (1991), Schmidt, Dick, Forbes, Tasker, eds., Elsevier Science Publishing,
451-454, 1992.
2. D. E. Grady and R. L. Moody, Shock Compression Profiles in Ceramics,
Sandia National Laboratories Report, SAND96-0551, March 1996
3. J. R. Asay and L. C. Chhabildas, Determination of Shear Strength of
Shock-Compressed 6061-T6 Aluminum, Shock Waves and High-Strain-Rate
Phenomena in Metals, Myers and Murr, eds., Plenum Pub. Corp, New York, NY
(1981)
4. L.C. Chhabildas, J.R. Asay, L.M. Barker, Shear strength of Tungsten
Under Shock and Quasi-Isentropic Loading to 250 GPa, Sandia National
Laboratories Report, SAND88-0306, April, 1988.
5. L. M. Barker and R. E. Hollenbach, Laser Interferometer for Measuring
High Velocities of any Reflecting Surface, J. Appl. Phys. 43, (1972), pp. 4669-
4675.
6. L. M. Barker, R. E. Hollenbach, Shock-Wave Studies of PMMA, Fused
Silica, and Sapphire, J. of Applied Physics, Vol. 41, No. 10, 4208-4226,
September 1970.
7. Acoustic velocity measurements on a-axis single crystal sapphire
performed by J. H. Gieske, Sandia National Laboratories, Albuquerque, NM.
246 Ceramic Armor Materials by Design
8. T. F. Thornhill, W. D. Reinhart, C. H. Konrad, L. C. Chhabildas, Accurate
Velocity Measurements of the Two-Stage Gun Projectile, 51st Aeroballistics
Range Association Meeting, September 17-21, 2000 Madrid, Spain.
9. J. L. Wise, L. C. Chhabildas, Laser Interferometer Measurements of
Refractive Index in Shock-Compressed Materials, Shock Waves in Condensed
Matter, Gupta, eds., (Plenum, New York), 441, 1986.
10. D. E. Grady, Dynamic Properties of Ceramic Materials, Sandia National
Laboratories Report, SAND88-3266, February 1995.
11. W. J. Carter, Hugoniot Equation of State of Some Alkali Halides, High
Temperature-High Pressure. 5:313 (1973)
12. T. Mashimo, Y. Hanaoka, K. Nagayama, Elastoplastic Properties under
Shock Compression of Al203 Single Crystal and Polycrystal, J. Appl. Physics. 63,
327 (1988)
13. R. A. Graham, W. P. Brooks, Shock-Wave Compression of Sapphire from
15 to 420 Kbar. The Effects of Large Anisotropic Compression., J. Phys. Chem.
Solids, Vol. 32, pps. 2311-2330 (1971), printed in Great Britain.
14. M. N. Pavlovskii, Shock Compression of Six Hard Substances, Soviet
Phys. Solid State, 12, 1736 (1971)
15. Y. Sato, S. Akimoto, Hydrostatic Compression of four corundum-type
compounds: Al203, V2O3, Cr2O3, and Fe2O3, J. Appl Phys., 50(8), August 1979.
16. M. D. Furnish, L.C. Chhabildas, Alumina Strength Degradation in the
Elastic Regime, Shock Compression of Condensed Matter, (1997), Schmidt,
Dandekar, Forbes, eds., pp. 501-504, 1997.
17. L. E. Pope, A. L. Stevens, Wave Propagation in Beryllium Single Crystals,
Metallurgical Effects at High Strain Rates, Rohde, Butcher, Holland, and Karnes,
eds., pp. 350-366 (1973).
18. G. I. Kanel, S. V. Razorenov, A. V. Utkin, V. E. Fortov, K. Baumung, H.
U. Karow, D. Rusch, and V. Licht, Spall Strength of Molybdenum Single Crystals,
J. Appl. Phys. 74 (12), December 1993.
19. G. R. Fowles, Dynamic Compression of Quartz, Journal of Geophysics
Res., 72, 5729, (1967).
Ceramic Armor Materials by Design 247
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COMPRESSIBILITY AND SHEAR STRENGTH OF TITANIUM DIBORIDE
UNDER PLANE SHOCK WAVE LOADING
D. P. Dandekar and E. J. Rapacki
U.S. Army Research Laboratory
AMSRL-WM-TD
Aberdeen Proving Ground, MD 21005-5066
ABSTRACT
Compressibility and shear strength of ceramics influence the potential
usefulness of these materials in protective systems against ballistic impact threats.
These properties of ceramics, due to their dominant brittle nature, can undergo
drastic changes under impact induced stress waves, and thus determine its
ultimate impact worthiness. This work brings together the results of shock wave
investigations on titanium diboride (TiB2) having a bearing on its compressibility
and shear strength. Further, the results of a limited number of shock–release–re-
shock (double-shock) experiments performed on TiB2 indicate that the observed
work–hardening response of TiB2 above the Hugoniot Elastic Limit (HEL) under
the initial shock compression augments its shear strength under the subsequent
shock wave loading. The results of this type of experiment are crucial for
developing a better understanding of the performance of ceramics under ballistic
impact, and for aiding the development of material models to predict the ballistic
performance of ceramics in armor configurations.
INTRODUCTION
Ceramic materials have received considerable attention from the ballistic
impact protection community for armor applications ranging from personnel
protection (body armor) to combat vehicle protection (integral composite armor).
Their very high strength to mass density ratio makes these materials attractive for
high performance, low weight armor system components. The asymmetry of the
compressive and tensile strengths of most ceramics, however, creates challenges
for their implementation. The compressibility and shear strength of a material
determine its dilatational and deviatoric deformation response to loading.
Because ballistic impact conditions produce propagating shock waves, which
subsequently are reflected from impedance mis-matched material interfaces, the
Ceramic Armor Materials by Design 249
To the extent authorized under the laws of the United States of America, all copyright interests in this publication are the propertyof The American Ceramic Society. Any duplication, reproduction, or republication of this publication or any part thereof, withoutthe express written consent of The American Ceramic Society or fee paid to the Copyright Clearance Center, is prohibited.
shock-release characteristics of armor materials must be known. The degradation
of strengths due to these load-unload conditions also provides insight for the
temporal longevity of the initial high strength characteristics. The shock-release-
reshock, or double-shock impact experiment1, 2, 3, 4
additionally provides a precise
methodology to probe a material’s retained strength, and hence its utility as an
armor system component, since ballistic loading times are typically of significant
duration.
This paper explains in some detail the rationale, implementation and analysis
of such double-shock experiments. A well-known high performance armor
ceramic, titanium diboride (TiB2) is the material that has been focused on in the
analysis of extant data, and the material investigated by the multi-shock impact
technique. The retained shear strength and work-hardening behavior observed in
this material’s response to plane shock wave loading helps to explain its excellent
performance as ballistic armor.
EXPERIMENTAL TECHNIQUE, ANALYSIS AND RESULTS
A schematic representation of double shock experiments is shown in Fig. 1.
Dandekar, Gaeta and Horie2 and Dandekar
3, 4 give detailed descriptions of the
double shock experiments. Briefly, an initially shocked material undergoes a
totally stress-free state for a pre-determined time duration before being shock
compressed a second time. There are two general configurations of this type of
experiment. In one (Fig. 1a), a thick impactor simultaneously impacts two
targets. One target consists of a single specimen, and the other consists of two
specimens separated by a measured, uniform gap between them. In the latter, the
firstly impacted target becomes the impactor of the second after stress release.
The free surface velocities are measured by multibeam interferometry, using
VISAR, (velocity interferometer system for any reflector). Alternatively, (Fig.
1b) two impactors, which are separated by a measured, uniform gap between
them, impact a target with an in-material gage to measure stress or particle
velocity in the target. Variables in both experimental configurations include: the
relative thicknesses of impactors and targets, the gap between the two targets or
two impactors, and the relative mechanical impedances of the impactors and
targets. The experimentally measured variables are: impact velocity, tilt of
impact, successive shock and release wave speeds, free surface velocity profiles
in the configuration of Fig. 1a, and stress wave profiles at various target locations
in the configuration of Fig. 1b. In either configuration shown in Fig. 1, a
transparent window could be used to monitor the particle velocity profile or stress
wave profile at the target-window interface. It should be mentioned that the
configuration shown in Fig. 1a does not permit observation of the change in wave
profile due to propagation of the second shock in a material, except in transparent
materials. The advantage of using an interferometer for inferring the shock
250 Ceramic Armor Materials by Design
properties of a material under multiple shocks and release is that it does not
require any calibration. The analysis of in-material wave profiles, obtained from
the configuration shown in Fig. 1(b), is straightforward.
Figure 1. Schematic configurations of shock, total release, and re-shock
experiments; (a) single impactor/multiple targets, whose free surface velocities
are monitored by multiple VISAR beams, and (b) multiple, successive
impactors/single target, which uses an in-situ stress or particle velocity gage to
monitor the wave profile.
Shock wave impact experiments were performed with a ten (10) cm diameter
single-stage gas gun at the U. S. Army Research Laboratory. The circular disc
specimens of TiB2 were such as to satisfy the one dimensional strain condition for
the total duration of wave profile measurements.4 The 32-50 mm diameter
ceramic disks were lapped flat to 5 m, and their opposing faces were mutually
parallel to within two (2) parts in 104 over the lateral dimension of the disks. The
deviation of planarity of impact in any given experiment was around 0.5 mrad.
Impact velocity of a projectile was determined by measuring the time intervals
signaled by the electrical shorting of four charged pins of known separation
distance. The precision of impact velocity determinations was 0.5%. The wave
Ceramic Armor Materials by Design 251
profiles were measured by means of a multibeam VISAR, or by means of a
Manganin stress gage. The precision of measurements of the particle velocity by
using VISAR, and of the stress measurements by means of Manganin gages, were
1% and 2.5%, respectively. The Manganin stress gages (Micro-Measurements,
Inc., Type LM-SS-125CH-048) are calibrated to nine (9) GPa.3
As mentioned above, the configuration of a repeated shock experiment shown
in Fig. 1(a) does not permit direct observation of the transmitted wave profile in a
material subjected to the repeated shock. However, the response can be
monitored through the material being subjected to the first shock and release.
Hence, the analysis is done in two stages. First, the response of the material
subjected to the first (initial) shock, generated by the impact of an impactor with
velocity v, and total release from this shock, are obtained as in a normal
transmission shock wave experiment. Incidentally, this can be conducted
simultaneously with the multiple shock experiments, as shown in Fig. 1(a), thus
insuring identical impact velocity for both experiments. In the second stage, the
analysis follows the procedure adopted for a front surface impact experiment, in
which the stress state attained during the second (subsequent) impact of the pre-
shocked material is obtained from the measured free surface wave profile, and the
measured response of the material during the first shock and its release.
Dandekar, in Ref. 4, presents the specific details of the analysis.
First Shock Response
Ambient properties and elastic compression: The ambient properties of the
TiB2 from various sources used in the various investigations to determine their
shock response are given in Table I. The two values of the Hugoniot Elastic
Limit (HEL) given in this table require some explanation. Plane shock wave
investigations on TiB2, irrespective of source and/or chemical impurities, show
the presence of two cusps prior to the onset of inelastic deformation. These two
cusps indicated two precursor waves, each propagating with velocity very nearly
the same as the elastic waves, but differing in magnitudes. The values of the first
and the second cusps were found to be between 4.2 - 5.9 GPa, and 8.0 - 17.0 GPa,
respectively. The first cusp was shown to be the limit of elastic deformation in
TiB2 in the sense that the material was damaged irreversibly above this first cusp.5
Compressive strength: Compressive strength or compressibility of a material
at different pressures is obtained from hydrodynamic compression of a material
under dynamic loading. Hydrodynamic compression of a material may be
represented by a functional dependence of the bulk modulus (K) on pressure. It is
well known that hydrodynamic compression of a hard material like a ceramic is
satisfactorily represented by its initial bulk modulus K0 and its initial pressure
derivative i.e., K0'. Thus, the values of K0 and K0' jointly convey the magnitude of
the compressive strength of a material. Dandekar and Benfanti6 analyzed the
252 Ceramic Armor Materials by Design
existing data on shock compression of TiB2. The study of the pressure
dependence of the elastic constants of TiB2 by Frankel, Abbate and Dandekar7
yields a value of 2.02 0.18 for the pressure derivative of the bulk modulus, K0'.
This value of K0' is in agreement with the high pressure shock data reported by
Gust, Holt and Royce8 and Marsh
9. This value of K0' together with the value of
bulk modulus K0 given in Table 1 for TiB2 yield its compressive strength.
Table I. Properties of various source TiB2 ceramic at ambient pressure condition.
Manufacturer Ceradyne
Eagle-
Picher Cercom
Union
Carbide (unknown)
Data Reference [6,7] [10] [10] [8] [11]
Mass Density,
(Mg/m3) 4.49±0.01 4.45 4.51 4.515±0.002 4.36±0.03
Elastic Wave
Velocities (km/s):
longitudinal, CL 11.23±0.11 10.93 10.79 11.21±0.20 10.79±0.15
shear, CS 7.41±0.13 7.30 7.43 7.30±0.16 7.24±0.10
bulk, CB 7.27±0.24 6.96 6.54 7.39±0.37 6.82±0.25
Elastic Constants:
bulk modulus, K0
(GPa)
237±16 216 193 247±12 203±15
Shear modulus,
(GPa)
246 ±9 237 249 241 ±5 228 ±6
Poisson’s ratio, 0.114±0.011 0.097 0.049 0.131±0.012 0.090±.009
HEL stresses, HEL:
1st cusp (GPa) 5.9 4.7 - 5.2 - - 4.2 - 4.9
2nd cusp (GPa) 13.5 13.1-13.7 17.0 8.0 9.0
Shear strength: The stress offset between the hydrodynamic compression and
the shock hugoniot data permits calculations of the shear strength of a material.
Shear strength for Cercom TiB2, from the shock hugoniot measurements to 60
GPa reported by Grady10
was calculated in this manner and reported in Ref. 6.
The values of shear strength of the Cercom produced TiB2, obtained from
simultaneous measurements of longitudinal and lateral stress under plane shock
wave loading12, 13
are compared with the those obtained on the same material from
the offset between the shock Hugoniot and the hydrodynamic compression of
TiB2.6, 7, 10
The values of shear strength of the TiB2 used in the investigation by
Winkler and Stilp11
are included for completeness. The values of shear strength
retained, and lateral stresses developed, in the various TiB2 under the first shock
wave propagation are given in Table II. The values of shear strength, , and
lateral stress, Y, given in parentheses in Table II are calculated from linear
Ceramic Armor Materials by Design 253
elasticity theory; i.e., Eqns. 1 and 2. In these equations X is the longitudinal, or
impact stress, and is Poisson’s ratio.
Table II. Shear strength and lateral stress imposed under plane shock wave
loading of TiB2; the materials are from the various sources indicated.
Compression
V/V0 (-)
Impact Stress
X (GPa)
Shear Strength
(GPa)
Lateral Stress
Y (GPa)
Eagle-Picher [Ref. 6 & 10]
0.9173 31.4 8.4 14.6
0.8929 46.7 14.6 17.5
Cercom [Ref. 6 & 10]
0.9825 8.6 3.8 (4.1)* 0.8 (0.4)*
0.9714 15.0 6.9 (7.1)* 1.2 (0.8)*
0.9428 24.5 9.3 5.9
0.9216 32.6 11.7 9.2
0.9207 32.2 11.2 9.8
0.8806 49.8 16.5 16.8
0.8635 61.0 21.2 18.6
Cercom [Ref. 12]
0.9876 6.8 3.2 0.4
0.9820 10.0 4.6 0.8
0.9615 19.5 8.4 2.7
0.9605 19.5 8.2 3.1
0.9474 24.0 9.2 5.6
Cercom [Ref. 13]
- 7.1 3.2 {2.5}** 0.7 {2.1}**
- 16.5 6.4 {4.5}** 3.7 {7.5}**
- 18.6 7.0 {6.0}** 4.6 {6.6}**
Unknown [Ref. 11]
0.9912 4.4 1.9 (2.0)* 0.6 (0.4)*
0.9800 9.1 3.7 (4.1)* 1.7 (0.9)*
0.9711 10.6 3.4 3.9
0.9678 14.3 5.6 3.1
Ceradyne [Ref. 3]
0.9846 8.7 3.7 (3.8)* 1.3 (1.1)*
0.9774 12.8 5.4 (5.6)* 2.0 (1.6)*
0.9545 19.1 5.6 7.9
0.9523 19.6 5.6 8.4
254 Ceramic Armor Materials by Design
0.9532 19.5 5.7 8.1
*Parameter by elasticity calculation **parameter for failed material
= (1-2 ) X /2(1- ) (1)
Y = X /(1- ) (2)
These values are comparable to those obtained from the stress offset between the
shock hugoniot and hydrodynamic compression of TiB2. These data indicate that
the shear strength retained by the various source TiB2 increases with an increase
in the magnitude of the impact stress under plane shock loading. This is most
clearly evident in Cercom and Eagle-Picher material because shock experiments
in these materials were performed at impact stresses exceeding twice the
magnitude of their respective HEL’s; see Table I. Further, Bourne, Gray and
Millet13
measured lateral stresses at two (2) mm and six (6) mm from the impact
surface in TiB2 specimens, and their lateral stress profiles show a two-step
structure. They associate the first step with the shear strength of intact material,
and the second step with the shear strength of damaged material, respectively.
They assumed that the damaged material was generated by the propagation of a
failure front in the material. The values of shear strength and lateral stress
associated with failed material are given in curly brackets in Table II. Figure 2
shows a plot of shear strengths of various TiB2 as a function of impact stress.
Second Shock Response
Dandekar4 performed the double shock experiments on Ceradyne material
only, to support then ongoing work on that material at the Army Research
Laboratory. The primary purpose was to examine whether the work hardening
behavior exhibited by Cercom and Eagle-Picher TiB2 was also exhibited by
Ceradyne TiB2 material. An analysis of those single- and double-shock
experiments on Ceradyne material at ~19 GPa is given in Table III. The data
show that the first shock of magnitude ~19 GPa, and subsequent release there-
from, follows an elastic-plastic (work-hardening) deformation path. The
subsequent second shock of ~19 GPa in this material is attained through elastic
deformation, because the impedance magnitude for the second shock is 55
Gg/m2s, that is, equal to its elastic impedance. As a consequence, the shear
strength of this TiB2 increases from 5.6 GPa under the first shock, to 8.3 8.7
GPa under the second shock. Further, the estimates of lateral stress imposed on
this material during the first and second shock decrease from 8.0 GPa to 1.7 2.0
GPa. The values of shear strength and lateral stress given in parentheses in Table
III were obtained using the elastic relations, Eqns. 1 and 2, and assumed that the
Poisson’s ratio remained unaltered, i.e., equal to 0.114. The shear strength
Ceramic Armor Materials by Design 255
value of 8.7 GPa and the lateral stress value of 1.7 GPa were obtained under the
assumption that the equation of state of an intact TiB2 is invariant. Using the new
value of shear strength , or lateral stress Y, in conjunction with the axial
compressive stress X, the new computed value of is 0.082. It is well known
that the Poisson’s ratio value of a solid does change under compression, but
whether the above-calculated new value is valid requires independent verification.
TiB2
0
5
10
15
20
0 10 20 30 40 50 60
Impact Stress, X (GPa)
Sh
ea
r S
tren
gth
, (
GP
a)
Eagle Picher [6,10]
Cercom [6,10]
Cercom [12]
Cercom [13]
Cercom [13]
(w/failure)Unknown [11]
Ceradyne [3]
Figure 2. Shear strength versus impact stress for TiB2 ceramics; the materials are
from the various sources indicated.
DISCUSSION OF RESULTS
The results of shock wave experiments indicate that titanium diboride exhibits
increasing shear strength with increase in impact stress. This behavior persists in
Cercom and Eagle-Picher materials to 46 and 61 GPa, respectively. The observed
work-hardening behaviour of titanium diborides is substantiated through the
results of a few shock-release and reshock experients in Ceradyne material. This
material deforms in an elastic-plastic manner under a first shock of magnitude
19.1-19.5 GPa, and maintains a shear strength of magnitude 5.6 GPa. When this
material is subjected to a second shock of the same magnitude following a
complete release from the first shock, it deforms elastically to 19 GPa. The value
of the shear strength under the second shock compression, calculated from the
offset between the shock hugoniot and the hydrodynamic compression, is 8.7
GPa. An estimate of the shear strength, using the elastic relation, Eqn. 1, with =
256 Ceramic Armor Materials by Design
0.114, is 8.3 GPa. These two values for the shear strength are within the
uncertainty of the measurements.
Table III. Summary of 1st and 2nd shock response of Ceradyne TiB2 at ~19 GPa.
Experiment #: 403 406 422
1st shock:
Axial stress, X (GPa) 19.6 19.5 19.1
Particle velocity, u (km/s) 0.448 0.445 0.433
Mass density, (Mg/m3) 4.714 4.710 4.704
Shear strength, (GPa) 5.6 5.7 5.6
Lateral stress, Y (GPa) 8.4 8.1 7.9
1st release:
Free surface vel., u (km/s) 0.803 0.837 0.781
Impedance, Z (Gg/m2s) 55 50 55
Mass density, (Mg/m3) 4.571 4.554 4.564
2nd shock:
Axial stress, X (GPa) - - 19.0
Particle velocity, u (km/s) - - 0.348
Impedance, Z (Gg/m2s) - - 55.
Mass density, (Mg/m3) - - 4.700
Shear strength, (GPa) - - 8.7 (8.3)*
Lateral stress, Y (GPa) - - 1.7 (2.5)*
* see DISCUSSION text
Bourne, et al.13
observed two steps in the lateral stress profiles of Cercom
TiB2 when shocked to between 7 and 19 GPa. Figure 2 shows that the shear
strengths based on the magnitude of the first step are in reasonable agreement
with those measured by Rosenberg et al.12
and reported by Dandekar and
Benfanti6. The magnitudes of shear strength at impact stresses from 7 to 19 GPa
decrease by 16% to 28% from their respective initial values due to the second
steps in the lateral stress wave profiles, see Table II. Bourne, et al.13
attributed
this reduction of shear strength to the subsequent propagation of a failure front,
which brings about a degradation of the shear strength. However, Murray and
Proud14
showed that the observation of two-step lateral stress profiles in ceramics
is dependent upon the geometry of the experimental configuration at a given
impact stress. Thus, the observed reduction in the shear strength of TiB2 could be
simply the manifestation of the geometrical configuration of the experiments in
Ref. 13. Further, the existence of failure front propagation in a solid under plane
shock wave compression is easily verified by the presence of recompression in the
Ceramic Armor Materials by Design 257
longitudinal wave profile. Therefore, such experiments must be performed on
TiB2 to independently verify the propagation of a failure front in the material, and
to corroborate the suggested degradation of the shear strength.
REFERENCES1D. Yaziv, S. Bless and Z. Rosenberg, “Study of Spall and Recompaction
Using a Double-Impact Technique”, Journal of Applied Physics, 58 [9], 3415-
3418 (1985). 2D. P. Dandekar, P. J. Gaeta and Y. Horie, “Double Shock and Release
Experiments in PMMA and Z-cut Sapphire,” pp. 281-284 in Shock Waves in
Condensed Matter - 1987, Edited by S. C. Schmidt and N. C. Holms, North-
Holland Press, New York, 1988.3D. P. Dandekar, “Response of Ceramics Under Single and Repeated Plane
Shock Wave Loading - A Case Study of Titanium Diboride,” pp. 242-253 in
Proceedings of IUTAM Symposium on Impact Dynamics, Edited by Z. Zemin,
Peking University Press, Beijing, PRC, 1994.4D. P. Dandekar, “Response of Protective Ceramics Under Single and
Multiple Impacts,” pp. 133-141 in Wave Propagation and Emerging
Technologies, AMD- 188, Edited by V. K. Kinra, R. J. Clifton and G. C. Johnson,
ASME Press, New York, 1994. 5D. P. Dandekar, “Effect of Shock-Re-Shock on Spallation of Titanium
Diboride,” pp. 487-490 in Shock Compression of Condensed Matter - 1991,
Edited by S. C. Schmidt, R. D. Dick, J. W. Forbes and D. G. Tasker, North-
Holland Press, New York, 1992. 6D. P. Dandekar and D. C. Benfanti, “Strength of Titanium Diboride Under
Shock Wave Loading,” Journal of Applied Physics, 73 [2], 673-679 (1993).7J. Frankel, A. Abbate and D. P. Dandekar, “Pressure Dependence of the
Elastic Constants of Polycrystalline Titanium Diboride,” pp. 881-884 in Recent
Trends in High Pressure Research, Edited by A. K. Singh, Oxford Press, New
Delhi, 1992. 8W. H. Gust, A. C. Holt and E. B. Royce, “Dynamic Yield, Compressional
and Elastic Parameters for Several Lightweight Intermetallic Compounds,”
Journal of Applied Physics, 44 [2] 550-560 (1973).9S. P. Marsh, p. 354 in LASL Shock Hugoniot Data, Edited by S. P. Marsh,
University of California Press, Berkeley, CA, 1980.10
D. E. Grady, “Dynamic Material Properties of Armor Ceramics,” Sandia
National Laboratories Report, SAND 91-0147, Albuquerque, NM, 1991. 11
W-D. Winkler and A. J. Stilp, “Pressure Induced Macro- and Micro-
mechanical Phenomena in Planar Impacted TiB2,” pp. 555-558 in Shock
Compression of Condensed Matter - 1991, Edited by S. C. Schmidt, R. D. Dick, J.
W. Forbes and D. G. Tasker, North-Holland Press, New York, 1992.
258 Ceramic Armor Materials by Design
12Z. Rosenberg, N. S. Brar and S. J. Bless, “Shear Strength of Titanium
Diboride Under Shock Loading Measured by Transverse Manganin Gages,” pp.
471-474 in Shock Compression of Condensed Matter - 1991, Edited by S. C.
Schmidt, R. D. Dick, J. W. Forbes and D. G. Tasker, North-Holland Press, New
York, 1992. 13
N. K. Bourne, G. T. Gray III and J. C. F. Millet, “On the Failure of Shocked
Titanium Diboride,” pp. 589-592 in Shock Compression of Condensed Matter -
1999, Edited by M. D. Furnish, L. C. Chhabildas and R. S. Hixson, American
Institute of Physics, New York, 2000. 14
N. H. Murray and W. G. Proud, “Measurement of Lateral Stress and Spall
Strength in Ceramics,” pp. 151-156 in Fundamental Issues and Applications of
Shock Wave and High Strain-rate Phenomena, Edited by K. P. Staudahammer, L.
E. Murr and M. A. Meyers, Elsevier Press, New York, 2001.
Ceramic Armor Materials by Design 259
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DYNAMIC INDENTATION DAMAGE OF CERAMICS
Do Kyung Kim, Chul-Seung Lee, and Young-Gu Kim
Dept. of Materials Science and Engineering
Korea Advanced Institute of Science and Technology
Taejon, Korea
Chang Wook Kim, and Soon Nam Chang
Agency for Defense Development,
Taejon, Korea
ABSTRACT
A modified Kolsky bar technique with a spherical indenter was applied to
evaluate the damage behavior of armor ceramics in dynamic indentation. Also,
a small explosive detonator was used for the dynamic indentation on ceramics.
In both experiments a bonded-interface specimen was useful to analyze the
subsurface damage after the concentrated dynamic loading on ceramics. A more
extensive quasi-plastic zone was observed in the dynamic indentation than in the
quasi-static loading. Microfracture behavior of damage zone in dynamic
indentation have almost the same features as those of quasi-statically damaged
ceramics.
INTRODUCTION
Ceramics have high hardness and elastic modulus, and these properties give
ceramics high wear and impact resistances. Some ceramic materials, such as
alumina, silicon carbide and boron carbide, are primary candidates for armor
applications.[1,2] However, the dynamic responses that relate to projectile
impact are not well understood, and it is the objective of the present study to
suggest one technique for characterizing the damage behavior during impact
loading, as a basis for identifying the material parameters that primarily influence
on dynamic impact. There were some reports on the crack evolution of brittle
ceramics during dynamic impact,[3-5] showing rate-dependent hardness. But few
studies have been reported on the damage evolution during dynamic impact
because of difficulties in recovering specimens after testing.
Indentation on the polished surface of a specimen with a spherical indenter,
Ceramic Armor Materials by Design 261
To the extent authorized under the laws of the United States of America, all copyright interests in this publication are the propertyof The American Ceramic Society. Any duplication, reproduction, or republication of this publication or any part thereof, withoutthe express written consent of The American Ceramic Society or fee paid to the Copyright Clearance Center, is prohibited.
Figure 1. Experimental set-up for the modified-Kolsky bar dynamic indentation
experiment.
which is generally called Hertzian indentation, have been extensively studied by
Lawn and colleagues to evaluate the properties of monolithic ceramics and
currently is being extended to layered structures.[6,7] In the course of analysis
of sphere-indentation, bonded-interface specimens could provide visualization
and quantitative analysis of damaged subsurface regions.
In this study, the modified Kolsky Bar technique with a spherical indenter and
a small explosive detonator, were used for dynamic indentation of typical armor
ceramics: alumina and silicon carbide. The subsurface damage zone during
dynamic indentation was characterized by using a bonded-interface specimen.
EXPERIMENTAL PROCEDURES
Specimen Preparation
Two armor ceramics, alumina(AD85, Coors Ceramics Co.) and silicon
carbide(hot-pressed SiC, Ceradyne Co.) were used for the experiment. To
reveal the subsurface damage, a special bonded-interface configuration was
used.[8] Polished surfaces of two half-specimens(6mm by 8mm by 35mm)
were glued face to face with a thin layer of adhesive under light clamping
pressure. Indentation was made with a tungsten carbide sphere across the
262 Ceramic Armor Materials by Design
Figure 2. Experimental set-up for the dynamic indentation by using small
detonator on (a) alumina and (b) silicon carbide ceramics.
interface trace. The two halves of the indented specimens were then separated
by dissolving the glue in acetone, cleaned, gold-coated, and examined by a
reflection optical microscope with Normalski interference illumination.
Dynamic indentation with modified Kolsky bar
To get the indentation with higher strain rate, compressed-gas driven Kolsky
bar equipment was used. The spherical indenter, tungsten carbide of 3.98 mm in
radius, was mounted on one side of a slender bar(100 mm by 10 mm and
placed on the interface of the bonded specimen. The other side of the slender
bar was impacted by the sabot-guided striker bar with a predetermined velocity,
which has the same length as the impact bar. The striker bar was accelerated by
a 20 mm-bore compressed-air gun. The impact was controlled by the velocity
of the striker bar, and the velocity was in the range of 5 to 15 m/s in the
experiments. Figure 1 shows the macroscopic view of experimental set-up.
Dynamic indentation with explosive detonator
With the same size of a bonded-interface specimen, the small explosive
detonator with a diameter of 5 mm was glued on the interface trace. Light vice
pressure was applied to avoid the shattering of the specimen during impact. The
separated two side of the specimen were cleaned, gold-coated, and examined by
the reflection optical microscope. Observation of the subsurface damage zone in
higher magnification was conducted by SEM.
Ceramic Armor Materials by Design 263
Figure 3. (a) Optical micrograph showing surface view(top) and side view
(bottom) of dynamically indented alumina. Indentation was performed in the
modified Kolsky bar set-up. The size of indenter was tungsten carbide of
3.98mm in radius and the striker velocity was 8 m/s. (b) Vickers hardness
variation as a function of radial distance from the center of top contact area.
Two data set at typical striker velocity are shown in (b). Vertical dashed lines
represent damage zone boundaries.
Hardness Measurement
Vickers hardness measurements have also been conducted on both the
damaged and undamaged area with load P =19.8 N. At least three indentations
were performed for each area. Hardness was determined as H = P/(2a2), where
P is applied load and a is impression half-diagonal. Application of Normarski
illumination enhanced the detection of surface impressions at each indent in the
optical microscope.
RESULTS AND DISCUSSIONS
Modified Kolsky bar experiment
Figure 3(a) shows optical micrographs of the top and side view of the
damaged alumina by the modified Kolsky bar. Impression and ring crack in top
surface and the extensive subsurface damage were observed. Figure 2(b) shows
Vickers hardness as a function of the radial distance from the center of contact
area at two typical striker velocities. It is clearly shown that the hardness of the
damage zone decreased as the distance decreased and the size of the damage zone
increased as the velocity increased.
264 Ceramic Armor Materials by Design
Figure 4. Side views of detonator-indented (a) alumina and (b) silicon carbide
ceramics showing quasi-plastic damage zone. Optical microscopy with Normalski
illumination highlights the detail contrast of damage zone.
Observation of Damage Zone
Figure 4 shows the subsurface side view of explosive detonator-indented (a)
alumina and (b) silicon carbide. To reveal the microstructure precisely, a mosaic
photo was made from each higher magnified micrographs. The damage zone
shows roughly semi-circular shape originated from contact area with the end of
explosive detonator. Damage scale of silicon carbide was smaller than that of
alumina. Lateral cracks were hypothesized to be generated by interference of
reflected pulses.
Ceramic Armor Materials by Design 265
Macroscopically, the "quasi-plastic" deformation zone developed in the strong
shear-compression region below the contact(the contact area is 5mm in diameter
in both specimen. The classical Hertzian cone crack developed from top surface
was hardly observed at the side view. The suppression of cone crack is
considered due to that the detonator develops only high shock pulse with
minimum mass of striking object.
Figure 5. SEM micrographs showing damage evolution from the detonator-
indented dynamic fracture. Microstructure of un-indented silicon carbide
with plasma etching reveals grain structure in (a). Higher magnification of
damage area in Figure 4(b) indicates that damage occurs at the grain boundary
of silicon carbide.
Higher magnification of damage zone in silicon carbide ceramics is shown in
figure 5(b) with comparison of the original un-damaged microstructure in (a).
Only microcracks in grain boundaries are clearly shown. In polycrystalline
ceramics, Lawn[9] has documented that the generic fracture mechanical model of
the microfracture evolution within the subsurface damage zone during a full
indentation loading and unloading. Microscopically, this indentation-induced
damage is associated with the activation of discrete "shear faults", from which
microcracks initiate. Interestingly the overall features of microfracture in
damage zone during dynamic indentation were almost same as that of quasi-
statically indented specimen.
Hardness of Damage Zone
Figure 6 shows the Vickers hardness values measured in damaged and un-
damaged zone of alumina and silicon carbide specimens that were indented by the
explosive detonator. Hardness of damage zone in alumina shows 43% of
original value, and that of silicon carbide shows 52% of original one. This
indicates that the damage severity of alumina is higher that that of silicon carbide.
It is considered that these hardness changes could be used as the indication of
damage severity during dynamic impact on ceramics.
266 Ceramic Armor Materials by Design
Figure 6. Vickers hardness of damaged and un-damaged region in alumina and
silicon carbide ceramics. Explosive detonator was used for the dynamic
indentation. Seven data points of radial hardness in damaged region were used to
calculate means and standard deviation.
CONCLUSIONS
Dynamic indentations on ceramics were introduced by the modified Kolsky
bar technique and the detonation of a small detonator. Subsurface damage zones
in alumina and silicon carbide ceramics were examined by using a special
bonded-interface technique. Described two techniques were suggested as a simple
and powerful technique to evaluate the damage response during dynamic impact
on ceramics. It is considered that the size and the hardness of damage zone can be
use to quantify the resistance of damage evolution during dynamic impact on
ceramics.
REFERENCES1M.L. Wilkins, C.F. Cline, and C.A. Honodel "Light Armor," UCRL-71817,
July 1969. 2R.C. Laible, Ballistic Materials and Penetration Mechanics, Edited by R.C.
Lable, Chapter 6 and 10, Elsevier Sci. Pub. Co, New York, NY. 1980. 3S.M. Wiederhorn and B.R. Lawn, "Strength Degradation of Glass Resulting
from Impact with Spheres," J. Am. Ceram. Soc, 60 [9-10] 451-58 (1977). 4A.G. Evans and T.R. Wilshaw, “Dynamic Solid Particle Damage in Brittle
Materials: An Appraisal, “ J. Mater. Sci., 12 [1] 97-116 (1977). 5D.B. Marshall, A.G. Evans, and Z. Nisenholz, “Measurement of Dynamic
Hardness by Controlled Sharp-Projectile Impact” J. Am. Ceram. Soc, 66 [8] 580-
Ceramic Armor Materials by Design 267
85 (1983). 6B.R. Lawn, "Indentation of Ceramics with Spheres: A Century after Hertz,"
Journal of the American Ceramic Society, 81 [8] 1977-94 (1998). 7K. S. Lee, S. K. Lee, B. R. Lawn, and D. K. Kim, Contact Damage and
Strength Degradation in Brittle/Quasi-Plastic Silicon Nitride Bilayers, Journal of
the American Ceramic Society, 81 [9] 2394-404 (1998). 8F. Guiberteau, N.P. Padture, H. Cai and B.R. Lawn, "Indentation Fatigue: A
Simple Cycle Hertzian Test for Measuring Damage Accumulation in
Polycrystalline Ceramics," Philos. Mag. A 69 [5] 1003-16 (1993). 9B.R. Lawn, N.P. Padture, F. Guiberteau, and H. Cai, "A Model for Microcrack
initiation and Propagation Beneath Hertzian Contacts in Polycrystalline
Ceramics," Acta Metall. Mater. 42 [5] 1683-93 (1994).
268 Ceramic Armor Materials by Design
TAYLOR-IMPACT EXPERIMENTS FOR BRITTLE CERAMIC
MATERIALS
L. C. Chhabildas and
W. D. Reinhart
Sandia National Laboratories
P. O. Box 5800
Albuquerque, NM 87185
D. P. Dandekar
Army Research Laboratory
Aberdeen Proving Ground, MD 21005-
5066
ABSTRACT
A new time-resolved test methodology is described which allows access to
loading rates that lie between split Hopkinson bar and shock-loading techniques.
Gas-gun experiments combined with velocity interferometry techniques have
been used to experimentally determine the intermediate strain-rate loading
behavior of Coors AD995 alumina, Cercom silicon-carbide and Cercom boron-
carbide rods. Graded-density materials have been used as impactors; thereby
eliminating the tension states generated by the radial stress components during the
loading phase. Results of these experiments demonstrate that the time-dependent
stress pulse generated during impact allows an efficient transition from the initial
uniaxial-strain loading to a uniaxial-stress state as the stress pulse propagates
through the rod. This allows access to intermediate loading rates over 5 x 103/s to
106/s.
INTRODUCTION
A new test methodology is described which allows access to loading rates that
lie between split Hopkinson bar and shock-loading techniques. Traditional split
Hopkinson bar techniques allow measurements on the failure stress of the
material at loading rates up to 103/s, where the definition of the failure stress is the
yield strength of the material determined under uniaxial-stress loading. In
contrast, plate impact techniques introduce uniaxial strain states at loading rates of
105/s or higher. At these very high strain rates the failure stress is defined as the
stress at which the material transitions from elastic deformation to plastic
deformation normally defined as the Hugoniot elastic limit of the material. The
experimental test methodology in each case prevents access to loading rates of the
order of 104/s. It is the purpose of this paper to report a new test methodology
Ceramic Armor Materials by Design 269
To the extent authorized under the laws of the United States of America, all copyright interests in this publication are the propertyof The American Ceramic Society. Any duplication, reproduction, or republication of this publication or any part thereof, withoutthe express written consent of The American Ceramic Society or fee paid to the Copyright Clearance Center, is prohibited.
that allows access to loading rates of 104/s. This is referred to, in this study, as
intermediate strain rate loading.
Taylor impact experiments consist of impacting a cylindrical rod onto a rigid
barrier [1]. Post-test observations or high-speed photography is then utilized to
determine the plastically deformed contour of the cylinder from which the
mechanical property data such as the dynamic yield stress also referred to as the
failure stress in this paper can be determined [2,3] through the measurements of
deceleration of the cylindrical rod. Most of the previous studies have been limited
to ductile specimens due to the ease with which the specimens can be recovered
for post-mortem analysis. In this test method, a rigid anvil is made to impact a
stationary sleeved-rod and its acceleration profile is used to estimate the dynamic
yield stress. The use of graded-density materials as a rigid anvil provides the time-
dependent loading profile.
A single-stage compressed gas gun combined with velocity interferometric
techniques [4] was used to experimentally determine the loading behavior of
ceramic rods. The rod dimensions are chosen so that the ratio of the length to its
diameter is at least four. Graded-density materials [5,6] were used to impact both
bare and sleeved ceramic rods [7-10] while the velocity interferometer [11] was
used to monitor the axial velocity of the free-end of the rods. This test-
methodology is well suited for brittle-ceramics because the ceramic rods will
invariably fracture during the loading process.
Results of these experiments demonstrate unique features of this novel test
methodology:
(a) a time-dependent stress pulse generated resulting from graded-density impact
allows a smooth and efficient transition from the initial uniaxial strain loading to a
uniaxial stress state as the stress pulse propagates through the rod,
(b) sleeved-rods in combination with graded-density impactors eliminate the
tension generated in the specimen during the loading stage.
(c) intermediate loading rates of 104/s obtained in this configuration lie in a
region which is not achieved easily by either split Hopkinson bar or shock-loading
techniques, and
(d) the loading rates can be varied from 104/s to 10
6/s through a combination of
increased impact velocity and different graded-density impactor design.
In this paper, only the results of experiments conducted on boron carbide rods
are reported. Results of Coors AD995 alumina rod experiments and silicon
carbide are published elsewhere [7-10].
270 Ceramic Armor Materials by Design
EXPERIMENTAL TECHNIQUE
These experiments were performed on a 64-mm diameter, smooth bore,
single-stage, compressed-gas gun that is capable of achieving a maximum
velocity of 1.2 km/s. Three electrically shorting pins were used to measure the
velocity of the projectile at impact. Four similar pins were mounted flush to the
impact plane and used to monitor the planarity of impact. Projectile velocity was
measured with an accuracy of about 0.5% and the deviation from planarity of
impact was about a milliradian. The graded-density impactor assembly is
fabricated by bonding a series of thin plates in order of increasing shock
impedance from the impact surface. The series of layered materials used in these
studies were TPX-plastic, aluminum, titanium, and 4340 steel. The thickness of
each layer is controlled to tailor the (time-dependent) input stress pulse into the
silicon carbide rod. The exact dimensions of each material assembly is given in
Table 1.
TABLE 1. Summary of impact experiments on sleeved silicon carbide Test
No.
Rod
Length/Diameter
(mm)/(mm)
Impactor
Materials
Impactor Thickness Impactor
Velocity
(km/s)
B4C-1 47.165/9.418 Steel/Ti/Al/TPX 12.73/0.224/0.244/0.244 0.309
B4C -2 47.061/9.416 Steel/Ti/Al/TPX 12.75/0.234/0.241/0.244 0.419
B4C-3 47.069/9.416 Steel/Ti/Al/TPX 12.52/0.234/0.249/0.251 0.600
B4C-4 47.089/9.416 Steel/Ti/Al/TPX 12.64/0.234/0.244/0.241 0.501
B4C-5 47.066/9.416 Steel/TPX 12.79/1.450 0.941
B4C-6 47.061/9.416 Steel/Ti/Al/TPX 12.71/0.328/0.243/0.509 0.945
B4C-7 47.628/9.418 Steel/Ti/Al/TPX 12.71/0.300/0.259/0.248 0.698
B4C-8 48.242/9.418 Steel/Ti/Al/TPX 12.66/0.300/0.260/0.268 0.799
This layered material assembly is used as facing on an aluminum projectile,
which is accelerated on a gas gun to velocities from 340 m/s to about 1000 m/s
prior to impact. This provides a time-dependent loading at the impact interface
from about 6 GPa to ~ 20 GPa, which is beyond the Hugoniot Elastic Limit for
the material [12-14]. The experimental target assemblies consisted of a sleeved
born carbide rod ~ 9.5 mm in diameter. The length of the rods in this study were
nominally 48 mm and 4340 steel was chosen for the close fitting sleeve material
to provide a good shock impedance to the boron-carbide sample. The outer
diameter of the sleeve was nominally 19 mm. The experimental configuration is
shown schematically in Fig. 1.
Ceramic Armor Materials by Design 271
TPX
Aluminum Ring
4340 Steel
TitaniumAluminum
Tilt Pins
VelocityPins
Sleeve:4340Steel
Boron Carbide
TungstenFoil
VISAR
Figure 1. Experimental configuration of a layered impactor and a ceramic rod target
assembly.
The boron carbide used in this study is obtained from Cercom. The density of the
material used in this investigation was 2.510 g/cm3; the longitudinal and shear wave
speed was determined to be 14.01 km/s and 8.83 km/s, respectively. This yields
9.60 km/s, 13.51 km/s, and 0.170 for the bulk-wave velocity, bar-wave velocity, and
Poisson’s ratio, respectively. Specifically, this is the same batch of material used in
previous studies on boron carbide [15] at the Army Research Laboratory.
A 0.033 mm thick tungsten reflector glued onto the free surface of the rod was used
to obtain the axial particle velocity measurements using the velocity interferometer,
VISAR having a time resolution of ~ 1 ns. The loading strain rate is varied either by
varying the impact velocity and/or by varying the thickness of the layered impactors at
the same impact velocity.
RESULTS
Eight experiments were conducted with boron carbide rods 9.5 mm in diameter and
48 mm in length with 4340-steel sleeves. The impact velocity was varied as shown in
Table 1, causing the stress and the loading rate to vary at the impact interface. Figure 2
shows the results of the experiments, B4C-1, B4C-2, B4C-3, B4C-5 and B4C-8, while
Figure 3 shows the results of experiments B4C-3, B4C-4, B4C-6, and B4C-7. The two
experiments at impact velocities below 0.6 km/s (B4C-1 and B4C-2, B4C-3, and B4C-4)
introduce stress levels that are at or below its Hugoniot elastic limit. The experiments
above 0.7 km/s are those above its Hugoniot elastic limit. The experiment B4C-3 which is
B4C
272 Ceramic Armor Materials by Design
at an impact velocity of 0.60 km/s exhibits a maximum free-surface velocity
measurement suggesting a stress at or approaching its Hugoniot Elastic limit. These
experiments are displayed in Figures 2 and 3, respectively.
0.941 km/s
0.419 km/s
0.309 km/s
0.799 km/s
0.60 km/s
0.0
0.1
0.2
0.3
0.4
0.5
0.6
0.7
3.3 3.5 3.7 3.9 4.1 4.3 4.5-0.1fr
ee-s
urf
ace
vel
oci
ty (
km
/s)
time (microsecond)
Figure 2. Free-surface velocity measurements at the free-end of the boron-carbide
rod resulting from graded-density impacts reported in Table 1. Experiments B4C-1,
B4C-2, B4C-3, B4C-5 and B4C-8 are shown in this figure.
Effect of Loading Rate
As indicated in Figures 2 and 3, the peak free-surface velocity measurements show an
increase with increased impact velocity up to an impact velocity of 0.6 km/s. At impact
velocities beyond 0.6 km/s the peak free-surface velocity measurement at the free-end of
the rod decreases with increasing impact velocity and is indicated in Figure 4. In Figure
5, the failure stress is plotted as a function of loading strain rate. A higher peak free-
surface velocity implies a higher yield stress also defined as the failure stress. This
provides experimental evidence for the dependence of failure stress upon loading rate.
There are reports of shear-strength loss in this material above its Hugoniot elastic limit in
shock experiments [12,13]. It should also be noted that the loading or the strain-rate also
increases with increasing impact velocity. Experiments B4C-5 and B4C-6 were performed
to investigate the effects of loading rate at the same loading stress. This was
accomplished by varying the dimensions of the layered impactor at the same impact
velocity (0.941 and 0.945 km/s, respectively).The free-surface velocity measurements are
comparable in these experiments, even though the loading rates differ by a factor of two –
the measurements suggesting that the failure strength has achieved its equilibrium value
of about 10.6 GPa .
Ceramic Armor Materials by Design 273
0.945km/s
0.501km/s
0.60km/s
0.698km/s
3.3 3.5 3.7 3.9 4.1 4.3 4.5-0.1
0.0
0.1
0.2
0.3
0.4
0.5
0.6
0.7
DISCUSSIONS
Previous studies on impact of ceramic rods have concentrated upon using a single
density impactor [16-18] to evaluate the uniaxial compressive behavior of the ceramics.
However, due to the low spall strength of ceramics [12-17] the radial stress components
will fracture the material during the loading phase, even though the mean stress of the
material indicates compression [16-18]. The technique proposed herein (i.e., using
graded-density impactors to study the uniaxial compressive behavior of the rods)
time (microsecond)
free
-su
rfa
ce v
elo
city
(k
m/s
)P
eak
fre
e-su
rfa
ce v
elo
city
, k
m/s
Impact velocity, km/s
Figure 3. Free-surface velocity measurements at the free-end of the boron carbide rod
resulting from graded-density impacts reported in Table 1. Experiments B4C-3, B4C-
4, B4C-6 and B4C-7 are shown in this figure.
0.3
0.4
0.5
0.6
0.7
0.8
0.9
1.0
0.3 0.4 0.5 0.6 0.7 0.8 0.9 1.0
Figure 4. Peak free-surface velocity measurements vs impact velocity for the profiles
show in figures 2 and 3.
274 Ceramic Armor Materials by Design
Grady
Lankford
Brar
2
4
6
8
10
12
14
16
10-4
10-3
10-2
10-1
100
101
102
103
104
105
106
Presentfa
ilu
re s
tres
s,
GP
a
strain rate, /s
Figure 5. Variation of failure strength with strain rate for boron carbide. Also shown are
results of quasi-static, Hopkinson bar and shock experiments.
Figure 6. Failure stress of Coors AD995 Alumina as a function of strainrate.
circumvents this problem by reducing the magnitude of tension generated in the ceramic
during loading [7]. A sleeved rod prevents the formation of radial tension during the
loading process [7]. The current experiments address strain rate effects in B4C at strain
rates of 104/s to over 10
5/s. This strain-rate regime is difficult to access either by
Hopkinson bar techniques or shock loading techniques. The leading edge of the axial
Ceramic Armor Materials by Design 275
compression wave traverses at an elastic-wave speed (14.0 km/s), followed by a second
compression wave traveling at a bar wave speed loading the material to a final stress at
strain-rates of ~ 104/s to 10
6/s. The first compression state l is calculated using 1=
( oc1 ufs)/2, where o is the initial density, c1 the elastic-wave speed, and ufs the
incremental free surface velocity measurement associated with the longitudinal elastic-
wave. The axial compression state a, and the loading strain rates d dt associated with
the bar wave are calculated using a= ( cb ufs)/2 and d dt = ufs/(2cbt), where cb is the
bar wave velocity, and ufs the corresponding free-surface velocity measurement, and t
the time duration for loading. Results of these experiments are shown in Figure 5, and are
compared to the low strain-rate Hopkinson bar experiments at strain-rates 10 3/s [19].
The low quasi-static strain-rate experiments yield a failure stress of ~ 5 GPa and shows
evidence of an increase to ~ 6 GPa at strain rates slightly above 103/s [19]. There is
considerable experimental scatter in the experiment suggesting the variability of the
material and perhaps the difficulty of conducting these experiments. As the strain rate
varies from 104/s to 6x10
5/s in these studies, the corresponding failure stress varies from
~ 6.5 GPa to a maximum of 11.1 GPa, before it approaches it’s equilibrium value of
10.6 GPa. This implies an Hugoniot Elastic Limit of at least 14 GPa for this material
before it sustains an equilibrium value of 13.3 GPa. The results clearly indicate the
dependence of the failure stress on the loading rates and also the loss in shear stress as
indicated in shock studies. The results lead credence to the hypothesis that damage
kinetics are rate-dependent, and ultimately, shock experiments yield higher estimates of
strength because rate-dependent kinetics prevent the nucleation and growth of
flaws/defects in materials during rapid loading.
The most significant result of this study is that the use of a graded-density impactor in
combination with sleeved rods allows accessibility to intermediate strain rates. Current
results for B4C, and previous studies on alumina and silicon carbide [4-5] both suggest
that the failure stress of ceramics is strain-rate-dependent. It should be noted that this
does not preclude the dependence of failure stress on mean pressure. It appears that
loading rates of a few times 104/s to 10
6/s can be achieved by optimizing the design of the
graded density layered materials, the diameter of the bar, and the impact velocity as
indicated in this investigation. One interesting study under consideration is to use the
graded-density materials as an impactor to perform isentropic loading experiments up to
its Hugoniot elastic limit. This will achieve lower loading rates than those obtained in
single shock experiments.
276 Ceramic Armor Materials by Design
2
4
6
8
10
12
14
1.E-
05
1.E-
04
1.E-
03
1.E-
02
1.E-
01
1.E
+00
1.E
+01
1.E
+02
1.E
+03
1.E
+04
1.E
+05
1.E
+06Strain Rate (S-1)
Steel Sleeved Rod - GDI
Steel Sleeved Rod - 1/2 GDI
Unsleeved Rod - GDI
Ta Sleeved Rod - GDI
Lankford
Grady, Feng, et al
106
102
100
10-2
10-4
104
Figure 7. Failure stress of Cercom SiC as a function of loading rate.
ACKNOWLEDGEMENTS
Sandia is a multiprogram laboratory operated by Sandia Corporation, a Lockheed
Martin Company, for the United States Department of Energy under Contract DE-AC04-
94AL85000. We acknowledge the able technical assistance provided by H. Anderson and
J. Martinez.
REFERENCES1G. I. Taylor, J. Inst. of Civil Engng. 26, pp. 486
2G. I. Taylor, Proc. R. Soc. London, Ser. A 194, pp. 289.
3J. C. Foster, Jr., M. Gilmore, L. L. Wilson, in Shock Compression of Condensed
Matter-2001, edited by M. D. Furnish, N. Thadhani, and Y. Horie, New York, AIP Press,
2002 (to be published). 4L. C. Chhabildas, and R. A. Graham, in Techniques and Theory of Stress
Measurements for Shock Wave Applications, 83, Edited by R. B. Stout, et. al., AMD,
1987 pp. 1-18. 5L. C. Chhabildas, L. N. Kmetyk, W. D. Reinhart, and C. A. Hall, Int. J. Impact
Engng. 17 (1995) pp. 183-194.6L. C. Chhabildas, J. E. Dunn, W. D. Reinhart, and J. M. Miller, Int. J. Impact Engng.
14 (1993) pp. 121-132. 7L. C. Chhabildas, M. D. Furnish, D. E. Grady, J. Phys IV FRANCE 7 (1997),
Colloque C3, (1997), pp. C3-137. 8L. C. Chhabildas, M. D. Furnish, W. D. Reinhart, D. E. Grady in Shock Compression
of Condensed Matter-1997, edited by S. C. Schmidt, D. P. Dandekar, and J. W. Forbes,
New York, AIP Press, 1998, pp. 505-508.
Ceramic Armor Materials by Design 277
9K. G. Holland, L. C. Chhabildas, W. D. Reinhart, M. D. Furnish, in Shock
Compression of Condensed Matter-1999, edited by M. D. Furnish, L. C. Chhabildas, and
R. S. Hixson, New York, AIP Press, 2000 pp. 585-588. 10
L. C. Chhabildas, W. D. Reinhart, Proceedings of the U.S. Army Symposium on
Solid Mechanics, edited by S. C. Chou and K. S. Iyer, (1999), pp 233-239. 11
L. M. Barker and R. E. Hollenbach, J. Appl. Phys. 43, (1972), pp. 4669-4675. 12
D. E. Grady, Dynamic Properties of Ceramic Materials, Sandia National
Laboratories Report, SAND94-3266, February 1995. 13
M. E. Kipp and D. E. Grady, Shock- Compression and Release in High-Strength
Ceramics, Sandia National Laboratories Report, SAND89-1461, February 1989. 14
N. S. Brar, Z. Rosenberg and S. J. Bless in Shock Compression of Condensed
Matter-1991, edited by Schmidt, S. C. and Dick, R. D., Forbes, J. W., Elsevier Science
Press, 1992, pp. 467-470. 15
D. P. Dandekar, Army Research Laboratory Report, ARL-TR-2456, April 2001. 16
A. Cosculluela, J. Cagnoux, F. Collombet, Journal de Physique IV, C3 (1991) pp.
109-116.17
N. S. Brar, and S. J. Bless in Shock-Wave and High-Strain-Rate Phenomena in
Materials, Edited by M. A. Meyers et. al., 1992, pp. 1041-1049. 18
J. L. Wise, D. E. Grady, High Pressure Science and Technology--1993, AIP
Conference Proceeding 309, Edited by S. C. Schmidt et. al., 1994, pp. 733-736. 19
J. Lankford, J. Amer. Ceramic Soc., 64, C33-C34 (1981)
278 Ceramic Armor Materials by Design
Analytical and Computational Modeling
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HISTORICAL PERSPECTIVE ON CERAMIC MATERIALS DAMAGE
MODELS
A.M. Rajendran
U.S. Army Research Laboratory
ARO, RTP, NC 27709-2211
ABSTRACT
Due to their high compressive strength, ceramic materials have been
frequently employed in armor systems for vehicle and soldier protection.
Ceramics are also candidate materials for ceramic engine components (turbine
blades, etc.) and aircraft engine containment systems due to their high temperature
properties. This paper presents a history of various approaches taken by
researchers to describe the brittle fracture of ceramics from the analytical
modeling of indentation processes to the recent high fidelity computational
modeling of projectile penetration processes in ceramic plates.
INTRODUCTION
Understanding and modeling of fracture in ceramic materials began with a
detailed study on quasi-static fractures induced by indentation loading. The
indentation modeling effort focused on the microcracking that occurs due to a
contact loading. During 1970's and 1980's, a large body of research work was
performed at various institutions and universities to characterize ceramic strength
through “hardness” measurements. Rajendran and Cook [1] presented a
comprehensive review of modeling of impact damage in ceramics. Lawn and
Wilshaw [2] reviewed the indentation fracture in detail.
Hockey [3] reported dislocation networks in alumina at local indentation sites.
Shockey et al [4] and Curran et al [5] address many of the deformation
mechanisms in confined ceramics under ballistic impact loading conditions.
These studies clearly established the presence of dislocations and twinning in the
brittle ceramics due to high pressures and high strain rate loading conditions.
Espinosa et al [6] reported evidence of inelastic deformation in compression due
to microcracking at triple junctions of the grain boundaries in recovered alumina
samples at velocities below the Hugoniot Elastic Limit (HEL). The
Ceramic Armor Materials by Design 281
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microplasticity in a brittle solid is often attributed to the dislocation motions in the
vicinity of microflaw tip regions. In brittle solids, large-scale grain distortions are
usually absent. Ewart and Dandekar [7] conducted a detailed microstructural
study of recovered titanium diboride specimens from low-velocity spall and
reshock experiments. Their study revealed that microcracks were the primary
form of irreversible damage in shock loaded ceramics. These microscopic
investigations indicated that the various forms of inelastic strain in the brittle
ceramics under shock and high strain rate of loading were caused by dislocations
and twins, microcracking, and pore collapse.
BACKGROUND
An accurate constitutive model must explicitly describe the various inelastic
processes through appropriate governing equations/laws. For instance, the axial
splitting and faulting in brittle material due to various levels of lateral
confinement were analytically modeled by Horii and Nemat-Nasser [8] based on a
wing-crack geometry under plane stress/plane strain loading conditions.
However, micromechanical modeling of the deformation processes under a three
dimensional stress/strain state is extremely difficult. Ravichandran and Subhash
[9] presented a micromechanical model for high strain rate behavior of ceramics
based on non-interacting microcracks that are uniformly distributed in the
material. This model was developed for biaxial compressive loading based on the
sliding crack (so-called wing crack). Hazell and Iremonger [10] reported a crack
softening damage model for ceramic impact and its application within a
hydrocode.
In summary, four distinct approaches can be identified as theoretical bases for
describing the inelastic deformation and fracture in ceramics. In the first
approach, the material is assumed to be elastic and stresses are calculated in the
finite element analysis based on Hooke’s law equations. Failure is predicted using
a generalized Griffith [11] fracture criterion. Mescall and Tracey [12] used the
Griffith criterion to model the fracture response of a ceramic armor in their HEMP
simulations. In the second approach, a numerical procedure is implemented in a
finite element/difference code to relax stresses to zero when some state variable
reaches a critical value. The ceramic behavior can be assumed to be either elastic
or elastic-plastic. Wilkins [13] implemented a fracture algorithm in the
Lagrangian finite-difference wavecode HEMP to examine a 0.30 CAL AP
projectile penetration into a thin ceramic plate backed by a metal substrate.
Recently, Anderson and Walker [14] adopted Wilkins' ceramic model to examine
ceramic dwell and defeat of the AP projectile.
In the third approach, the material is assumed to behave as an elastic-plastic
solid. This approach ignores the details of crack growth and concentrates on
282 Ceramic Armor Materials by Design
describing the effects of localized fracture on stress wave propagation. The
stiffness of the ceramic is not degraded due to damage, but the strength is
degraded due to plastic deformation induced damage. The concepts and equations
are the same as those derived for metal plasticity. Johnson and Holmquist [15]
modeled the strength of the ceramic as a function of pressure and strain rate
through a two-surface approach. Basically, there are two surfaces: one
corresponds to D = 0, and the other to D = 1. Steinberg [16] proposed a ceramic
damage model that is very similar to his metal fracture model. Rajendran and
Kroupa [17] modified the constitutive equations based on fragmentation [18] to
describe the shock response of silicon carbide. Recently Simha [19] proposed a
similar model.
In the fourth approach, fracture mechanics based microphysical theories and
models are employed to describe the deformation due to the compressive failure
processes. The basic idea of a microphysical model is to describe the apparent
inelastic behavior while keeping track of the microstructural evolutions under a
given set of loading conditions. For ceramics, the evolution laws for
microcracking must incorporate the fracture mechanics theories that describe the
conditions for crack growth. A statistical description of number of cracks per unit
volume as a function of position, crack size, and orientation, is an example of a
microphysical approach. The Hooke’s law based elasticity equations are
combined with an effective moduli description that relates the microstructure to
the macroscopic material properties. The models by Rajendran and Grove
[20,21], Addessio and Johnson [22], and Espinosa [23] follow this approach.
Until 1988, there were hardly any material models in hydrocodes (shock-wave
propagation based finite element/difference numerical codes) that could describe
the inelastic behaviors of brittle ceramics under shock and high strain rate loading
conditions. During the past decade, several new ceramic damage models (Partom
[24], Sadyrin, Ruzanov, and Podgornova [25], and Malaise, Tranchet, and
Collombet [26]) have been reported for hydrocode applications, each based on
one of these four approaches.
BRIEF SUMMARY OF A FEW CERAMIC MODELS
This section summarizes a few ceramic models that have been implemented
into various hydrocodes. Researchers around the world have proposed several
other models; no attempt is made to include them all in this brief summary.
Wilkins' Computational Scheme [13]
In a simplistic approach, Wilkins employed a two dimensional hydrocode
(HEMP) in a computational analysis of the impact and penetration of thin
laminate armor. He used the following simplified criteria: (1) fracture initiates on
Ceramic Armor Materials by Design 283
a surface, (2) a maximum principal stress greater than 0.3 GPa in tension causes
fracture, (3) there is a time delay for the complete fracture of a zone, (4) a
fractured zone becomes a source for the fracture of a neighboring zone, and (5)
fracture occurs only within a range of distance equal to or less than the time step
times the crack velocity in ceramic.
Figure 1. Effect of fracture strength on time for fracture (left column: a fracture
strength of 3 Kbar; right column is for 8 Kbar). Time is in microseconds.
284 Ceramic Armor Materials by Design
Using this numerical scheme, Wilkins modeled the evolution of the fracture
conoid in thin ceramic targets. The time delay for complete fracture is related to
the time for a crack to propagate across a computational cell. The crack speed is
assumed to be a fraction of the elastic shear wave speed. A value of 0.5 was
employed for the fraction. This implementation prevents cracks propagating from
cell to cell at an unrealistic speeds. He could successfully reproduce some of the
observed fracture patterns in the ceramic plate due a bullet penetration. Figure 1
illustrates the effects of fracture strength on time for fracture. Recently, Anderson
and Walker [14] successfully adapted Wilkins approach to model the dwell and
defeat of a 0.30-CAL AP projectile. This approach is limited to modeling thin
ceramic plates.
Liaw, Kobayashi, and Emery Computational Model [27]
A mechanically consistent model of impact damage based on elastic fractures
due to both tensile and shear loading is assumed in the simulation of dynamic
indentation for a spherical projectile on a structural ceramic. The projectile is not
explicitly modeled in the finite element modeling; instead, a transient contact load
is prescribed at the impact site. The impact is assumed to be elastic and the
surface loading conditions are derived from analytical solutions. The
implementation of their fracture algorithm includes the following steps: 1) a crack
will form at a material point perpendicular to the direction of the maximum
principal stress when this stress exceeds the ceramic’s tensile strength, 2) a set of
orthogonal cracks (parallel to the maximum shear direction) is assumed to form
when the maximum shear stress exceeds the shear strength of the ceramic, 3)
cracks are allowed to carry subsequent compressive loads according to Coulomb’s
law of dry friction, and 4) an element's stiffness across an open crack vanishes and
Figure 2. Damage patterns in a ceramic plate impacted by a hard (elastic) sphere.
Ceramic Armor Materials by Design 285
returns to its initial stiffness when the crack closes. Using this procedure, they
mapped the observed crack patterns (median crack, radial cracks, cone crack, and
lateral cracks regions) in the indented samples with reasonable accuracy as shown
in Figure 2. No attempt was made to employ this modeling approach to describe
shock wave profiles.
Modified TCK Model [17,18]
Rajendran and Kroupa [17] presented a modified version of the Taylor, Chen,
and Kuszmaul (TCK) [18] model that was developed to describe the brittle
behavior of oilshale under impact loading. In the modified version, the ceramic is
assumed to flow plastically under compression and no damage is allowed under
compression. In tension, the ceramic behaves in a brittle manner without any
plastic flow. A tensile damage parameter is defined using an expression that
combines the expressions derived by Kipp and Grady for fragmentation [28] and
Budiansky and O’Connell for a cracked solid [29]. The salient equations are
summarized as follows:
Compression: Y (1))D1()lnB1(Ys
Tension: ij (2)ijijkk e)D1(G2e)D1(K3
Damage is described by:
32
IC3dd
2
c
K20
2
1a;aNC;C
21
1
9
16D (3)
Ys is the static compressive strength, B is the strain rate sensitivity parameter, K
and G are bulk and shear moduli respectively, eij are deviatoric strains, ij are total
stresses, Cd is crack density, KIC is static fracture toughness, is degraded
Poisson's ratio, is a geometric factor, a is the crack (fragment) size, c is the
sound speed, N is the number of flaws, is material density, and is the strain
rate. This model was implemented in the EPIC code [30] and an experimentally
measured shock profile for silicon carbide was successfully reproduced.
Rajendran-Grove Model [20,21]
In this model, the total strain is decomposed into elastic ( ) and plastic ( )
strains. The elastic strain consists of the elastic strain of the intact matrix material
and the strain due to crack opening/sliding. Plastic flow is assumed to occur in
eij
pij
286 Ceramic Armor Materials by Design
the ceramic only under compressive loading when the applied pressure exceeds
the pressure at the Hugoniot elastic limit (HEL). Pore collapse during shock
loading is modeled using a pressure dependent yield function and the strains due
to pore collapse are assumed to be viscoplastic. The constitutive relationship for
the cracked material is given by:
(4)eklijklij M
The components of the stiffness tensor M are described by Rajendran [20]. The
microcrack damage is measured in terms of a dimensionless microcrack density .
The maximum microcrack size a is treated as an internal state variable.
Microcracks are assumed to extend when the stress state satisfies a generalized
Griffith criterion. This criterion requires the fracture toughness KIC as well as a
dynamic frictional coefficient as model parameters. The damage evolution law
is derived from a fracture mechanics based relationship for a single crack
propagation under dynamic loading conditions:
2n
I
crR1
G
G1Cna (5)
where CR is the Rayleigh wave speed, Gcr is the critical strain energy release rate
for microcrack growth, and GI is the applied strain energy release rate. The model
constants n1 and n2 are used to limit the microcrack growth rate. Under tension,
these two constants are assumed to be equal to one. The ceramic is assumed to
pulverize under compression when reaches a critical value of 0.75. The ceramic
model has six parameters to describe microcracking of the intact ceramic. This
model has been implemented in the EPIC code [30] and simulations of several
shock and impact configurations have been successfully performed.
Espinosa’s Multi-plane Model [23]
This model assumes that microcracking can occur on a discrete number of
orientations. Espinosa et al [23] selected nine orientations at intervals of 450
along three mutually perpendicular planes. The inelastic strain is entirely due to
(penny-shaped) microcrack opening/sliding of the cracks oriented normal to those
nine directions. The average inelastic strains are given by,
9
1k
)k(j
)k(i
)k(j
)k(i
)k()k(cij bnnb
2
1SN (6)
Ceramic Armor Materials by Design 287
The superscript k represents the orientation, N is the number of flaws per unit
volume, S denotes the surface of the microcrack, n is the corresponding unit
normal, and b is the average displacement jump vector across the surface S. The
b have been analytically derived for normal tractions under both tension and
compression. The corresponding expressions are:
)k(i
)k(l
)k(jjl
)k(jij
)k(2
)k(i nnnn2a
)2(E3
)1(16b (7)
and)k(
i)k(
2)k(
i fa)2(E3
)1(32b , (8)
where a(k)
is the crack radius of the penny-shaped microcracks on orientation k
and f (k)
is the effective shear traction vector on orientation k. The microcrack
growth law is very similar to the one that Rajendran and Grove [20,21] employed;
the multi-plane model uses the stress intensity factor instead of the strain energy
release rate. There are two crack growth constants, n1 and n2; these two constants
can take on different values for tension and compression.
Johnson-Holmquist Model [15]
There are two versions of this model: JH1 and JH2. The differences between
these two versions are very subtle. The JH model is a phenomenological model
based on an elastic-viscoplastic approach. The strength of the ceramic is assumed
to vary with pressure, strain rate and tensile strength. Basically, there are two
surfaces; one corresponds to D = 0, and the other to D = 1. Once damage initiates,
the flow surface reduces to an intermediate state, and at D = 1 the strength lies on
the second surface. As in the Johnson-Cook fracture model for metals, the
damage (D) increases with effective plastic strain. This model has been discussed
in detail elsewhere in this volume. Recently, several impact and penetration
configurations have been successfully modeled using the JH model.
Steinberg Model [16]
Steinberg assumed that all thermomechanical behaviors could be represented
through certain macroscopic variables, such as strain, strain rate, temperature, and
pressure. He adopted the salient features of his metal model for under high strain
rate and shock loading applications. The ceramic model assumes that both the
yield strength and shear modulus vary with respect to temperature and pressure.
However, the yield strength (Y) varies also with the strain rate. The governing
equations are:
288 Ceramic Armor Materials by Design
dT
dG
G
1Band
dP
dG
G
1A;)T(B
PA1G)T,P(
oo3/1oG , (9)
where is the compressive strain, is effective strain rate, and G, P, and T are
shear modulus, pressure, and temperature, respectively. Go is the shear modulus
of the ceramic before shock loading. The expression for the yield strength is:
n2ICoo
oA
n KC3D;G
GYDY , (10)
where o is the initial density of the ceramic, Co is the longitudinal wave speed,
KIC is the fracture toughness, n is a model parameter, and, according to Steinberg,
YA is a material constant which can be sample dependent, as it is a function of
purity, grain size, previous mechanical history, etc. The strain is simply
decomposed into elastic-viscoplastic, and the stress calculation in the
computational implementation follows conventional plasticity theories. Basically,
there are five model parameters: Go, YA, A, B, and D (or n). Steinberg also
employed a traditional void growth (spall) model to describe tensile cracking in
ceramics; the spall model requires three additional constants.
Addessio - Johnson Ceramic Model [22]
In this model, the inelastic strains in ceramics due to penny-shaped microcrack
growth under tension and compression are determined by integrating the
individual crack strains over a material volume, as well as all crack sizes and
orientations. By invoking several assumptions regarding the nature of crack size
and orientation, it is then possible to obtain simplistic expressions for the inelastic
strains. Addessio and Johnson [22] derived the following relationships for the
deviatoric parts of the inelastic strain components ( e ):cij
G
N
2
1
15
64where
)tension(6;)ncompressio()5(2,Sce
o
eeij
3ecij
(11)
In this equation, Sij are the stress deviators, c is the sound wave speed, is
Poisson’s ratio, and No is the number of flaws. In the deviatoric elastic stress-
strain relationship, the shear modulus is degraded through the following
expression:
Ceramic Armor Materials by Design 289
3e cG1
GG , (12)
where is a material model parameter that was arbitrarily introduced into the
above expression. Addessio and Johnson obtained crack growth criteria by
considering an energy balance on isolated cracks. The crack growth rate is
described by:
smax dtanhcac , (13)
where is assumed to be the shear wave speed, “a” is a factor that will
reduce the crack speed, and d
maxc
s is a measure of the distance the state of stress
exceeds the damage surface. The main model parameters are: c , No , o , , ,
and a. In addition to these parameters, the model also requires one or two other
parameters.
Riuo-Cottenot-Boussuge Tensile Damage model [31]
This model assumes that the ceramic deforms elastically below the Huguenot
Elastic Limit. When the shock amplitudes exceed the HEL, the ceramic deforms
plastically as metals. In the model, damage initiation occurs when the principal
stress exceeds a threshold stress th. When all the principal stresses are tensile
and exceed this initial threshold value, damage initiates in the planes that are
perpendicular to the principal stress directions. Therefore, damage can initiate
and propagate when . Note that the initial value ofcri cr is th. This
threshold stress is reduced according to: . The definition of dthicr d1 i is
the ratio of total crack surface of penny-shaped cracks over the total solid surface.
The corresponding expression for di is:
2i
si2
aNd . (14)
Ns and the initial microcrack size ai are model parameters. Since damage is
assumed to increase from 0 to 1, this assumption puts a bound on the maximum
crack size. In the model, the crack extends at a constant speed, Vf . This
extension is possible only when the initiation criterion as well as the time rate of
change of the principal stress is positive. The stresses are assumed to degrade
according to (1 - di ) e, where e is the principal elastic stress. A Mohr-Columb
290 Ceramic Armor Materials by Design
law is employed to describe the strength of the pulverized ceramic.
This model was used to simulate a three dimensional configuration in which a
steel cylinder (20 mm long, 11 mm in diameter) impacts a silicon carbide beam
(20 mm thick). In the experiments, stress measurements were made by
embedding a stress gauge between the back surface of the beam and a thin steel
plate, and photographs of the fracture patterns in the beam were obtained from a
high-speed camera. The authors validated their model through three-dimensional
simulations of this test configuration and compared the measured stress histories
with the computational results. The simulated crack patterns qualitatively agreed
with the observed crack patterns. The model also predicted the stress gauge
measurements well; however, an elastic analysis without damage also matched the
experiments reasonably well.
Simha’s Phenomenological Model [19]
Simha assumes that the ceramic fails at the Hugoniot Elastic Limit (HEL).
Microcracking due to sliding is assumed to be the dominant inelastic mechanism.
The Mohr-Coulomb law describes the strength of the failed ceramic. The
effective strength (Y) of the inelastic state is defined by,
)P(2
e3)P(YqsY , (15)
where P is the pressure, e is the effective deviatoric strain rate, is a parameter
that controls the contribution of the rate dependent term to the strength of the
failed material, and Yqs is defined as the rate independent part of the strength (like
a reference strength). Yqs is constant before the ceramic fails, and follows the
Mohr-Coulomb law (up to Ycap) after failure. Simha successfully used this model
to describe the shock response and penetration resistance of aluminum oxide.
VALIDATION AND VERIFICATION
The most frequently reported experimental technique to calibrate the high
strain rate and shock behaviors of ceramics is based on the plate impact test
configuration. In this configuration, a flyer plate is impacted against a target of
the same or different material at high velocity. The diagnostic measurements
include the use of either a peizo-resistive stress gauge or a velocity interferometry
system (VISAR). The measured wave profiles are often used in the calibration of
ceramic model constants. It is not possible to determine the exact nature of the
deformation processes from the measured profiles. However, microscopic studies
on the scientifically recovered targets often reveal many of the
Ceramic Armor Materials by Design 291
deformation/damage processes.
Grady and Wise [32] obtained particle velocity wave profiles (VISAR Data)
for various ceramic materials, including silicon carbide (SiC), boron carbide
(B4C), and titanium diboride (TiB2). The impact velocities in those experiments
were about 1500 and 2500 m/s. Most of the models reproduced the VISAR data
very well. For example, Figure 3 shows a comparison between the plate impact
data and the computed wave profiles using the model developed by Rajendran and
Grove [20,21] for four different ceramics.
Several other impact experimental configurations are available for model
validation. Attempts have been made to match the measured profiles from a wide
AD995 – ALUMINUM OXIDE (1.943 KM/SEC)
Time ( s)
0.5 1.0 1.5 2.0 2.5 3.0
Velo
cit
y (
km
/s)
0.0
0.2
0.4
0.6
0.8
1.0
1.2
1.4
ExperimentRG Model
S
Figure 3. Comparisons between simulations and data for four different ceramic materials.
292 Ceramic Armor Materials by Design
variety of experiments with one set of model constants. Holmquist, Templeton,
and Bishnoi [33], Simha [19], and Grove and Rajendran [34] have shown
successful comparisons between their model predictions and the data. At least
one or two parameters have to be calibrated to match the measured depth of
penetration (DOP) data from the projectile penetration experiments. Since the
governing equations in the model formulation are not precisely derived to describe
the actual micro/macro damage processes in a ceramic under a wide range of
stress / strain / strain rate, it is not surprising to discover that the model parameters
calibrated solely from the shock wave experimental data are not suitable for
predicting the DOP. For completion, brief descriptions of several other
experimental configurations are discussed in the following sub-sections.
Rod-on-Rod Impact
A short ceramic rod is made to impact a long stationary ceramic rod. In this
uniaxial stress configuration, a stress gauge is typically embedded in the target rod
to record the stress history. Fracture initiates at the impact end, with several
splitting type macrocracks forming and propagating toward the gauge location.
The measured peak stress from this experiment can be used in a qualitative sense
to validate the model constants under a uniaxial stress state.
Graded-Density Plate-on-Rod Impact
Recently, Chhabildas et al [35] reported an experimental configuration in
which a ceramic rod (L/D 4, sleeved or unsleeved) was impacted by a graded-
density flyer plate consisting of extremely thin (0.1-cm thick) layers of titanium,
aluminum, and TPX bonded to a 1.9-cm thick steel plate. A VISAR was used to
record the axial particle velocity of the free end of the target rod. This test
configuration generates a time-dependent stress pulse that smoothly and
efficiently transits from the initial uniaxial strain loading to a uniaxial stress state.
Also, the intermediate loading rates obtained in this configuration are not easily
achieved by either split Hopkinson bar or conventional shock-loading techniques.
Depth of Penetration Experiment
In the projectile penetration experiment, a tungsten long rod projectile is
launched at a nominal velocity of 1.5 km/s onto a thick square ceramic tile that is
laterally confined by a steel frame; the target assembly (tile and frame) is
mechanically clamped to a thick steel backup block. The depth of penetration
(DOP) of the projectile in the backup steel plate is measured and used as a
parameter to compare in the validation and verification of a model. High speed
photographs and X-ray radiographs are also often obtained as part of the
diagnostic measurements.
Ceramic Armor Materials by Design 293
SUMMARY
During the 1970's, the computational analysis of a projectile (metallic spheres
and cylinders) impacting a brittle ceramic plate was mainly performed to gain a
fundamental understanding of complex crack patterns developed due to the
impact. A combined indentation-based experimental and computational analysis
approach was employed in the evaluation of hardness and compressive strength of
ceramics. The response of the ceramic was assumed to be elastic in the
indentation analysis. During the past decade, researchers realized an urgent need
for a fully three-dimensional constitutive description of ceramic materials to
perform realistic hydrocode analyses suitable for impact-resistance applications.
Constitutive model formulations have mainly focused on incorporating the effects
of pressure, defects (pores and microcracks), and strain rate on strength and
stiffness of the ceramic. A few models have included the effects of flaw
orientation and/or microplasticity (dislocations, twins, etc.) on the degradation of
strength and stiffness. Those model parameters that cannot be directly measured
from experiments are estimated (calibrated) based on their ability to reproduce or
match the measured wave profiles. Most models use the Mohr-Coulomb law to
describe the compressive/shear loading response of the comminuted ceramic.
Generally, one or two parameters are needed for this purpose. Currently, there is
no physics-based model to accurately describe the comminuted ceramic response.
Curran et al [5] reported a micromechanical model based on the non-elastic
sliding and ride-up of fragments of comminuted particles. Their simulation of a
projectile penetration into a confined ceramic showed that the DOP is controlled
by 1) friction between the comminuted granules, 2) compressive strength of the
intact ceramic, and 3) compaction strength of the comminuted ceramic.
Current models are incapable of describing the effects of grain size and grain
boundary properties on the impact and shock resistance of ceramics. Recently,
Zavattieri and Espinosa [36] presented a grain level analysis of ceramic
microstructures subjected to impact loading. Through a two dimensional
stochastic finite element analysis, they explicitly modeled the details of grain
morphology and its effects on crack nucleation and propagation at grain
boundaries. Though we have made significant progress in modeling the ceramic
damage under shock and penetration loading conditions, there are still issues that
need attention and additional research.
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294 Ceramic Armor Materials by Design
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P.J. Hazell and M.J. Iremonger, “Crack Softening Damage Model for
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M.L. Wilkins, “Third Progress Report of Light Armor Program,” UCRL-
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C.E. Anderson and J.D. Walker, “Ceramic Dwell and Defeat of the 0.30-
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US Army Solid Mechanics
Symposium, Eds., K. Iyer and S.C. Chou, 17-28 (1999). 15
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16D.J. Steinberg, “Computer Studies of the Dynamic Strength of Ceramics,”
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Dimensional Stress Experiments and Constitutive Modeling,” Ph.D Thesis,
University of Texas at Austin, TX, 1998. 20
A.M. Rajendran, “Modeling the Impact Behavior of AD85 Ceramic Under
Multiaxial Loading,” Int. J. Impact Engng., 15 (6) 749-768 (1994). 21
A.M. Rajendran and D.J. Grove, “Modeling the Shock Response of Silicon
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H.D. Espinosa, P.D. Zavattieri, and S.K. Dwivedi, “A Finite Deformation
Continuum/Discrete Model for the Description of Fragmentation and Damage in
Brittle Materials,” J. Mech. and Phys Solids, 46 [10] 1909-1942 (1998). 24
Partom, “Calibrating a Strength Model for AD995 Alumina from Plate
Impact VISAR Profiles,” J. de Physique IV c8-495 (1994). 25
A. Sadyrin, A. Ruzanov, and T. Podgornova, “Modeling Impact Loading and
Failure of Brittle Solids; Constitutive Damage Model for AD995,” Research
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1121-1124 (2000). 27
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Damage in Structural Ceramics,” Journal of the American Ceramic Society, 67 [8]
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31P. Riou, C.E. Cottenot, and M. Boussauge, “Steel Rod Impact on Silicon
Carbide Beams: Experiments and Anisotropic Model of Damage,” in “Structures
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Computational Mechanics Publications, 533-541 (1996). 32
D.E. Grady and J.L. Wise, “Dynamic Properties of Ceramic Materials,”
Sandia Report, SAND93-0610, September (1993). 33
T.J. Holmquist, D. Templeton, and K. Bishnoi, “Constitutive Modeling of
Aluminum Nitride for Large Strain, High Strain Rate, and High Pressure
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Ceramic Failure Model,” a Chapter in PAC RIM IV proceedings, 2002. 35
L.C. Chhabildas, M.D. Furnish, W.D. Reinhart, and D.E. Grady, “Impact of
AD995 Alumina Rods,” Proceedings of the Shock Compression of Condensed
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508 (1997). 36
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Microstructures Subjected to Impact Loading,” a Chapter in PAC RIM IV
proceedings, 2002.
Ceramic Armor Materials by Design 297
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A COMPARISON OF CERAMIC MATERIAL MODELS
Douglas W. Templeton
U. S. Army Tank Automotive Research, Development, and Engineering Center
Warren, MI 48397-5000
Timothy J. Holmquist
Network Computing Services Inc./Army HPC Research Center
Minneapolis, MN 55415
Hubert W. Meyer, Jr., David J. Grove, and Brian Leavy
U.S. Army Research Laboratory
Aberdeen Proving Ground, MD 21005-5066
ABSTRACT
This paper presents results of a study investigating two different ceramic
models using two different computer codes and comparing their performance for
terminal ballistic problems. Computations were performed using the Johnson-
Holmquist (JH-1) and the Rajendran-Grove (RG) constitutive models for brittle
materials, for penetration problems into ceramics as implemented in both the
Eulerian CTH and the Lagrangian EPIC shock physics codes. The results of the
computations are compared to each other and experimental data, and an
assessment is made of the models’ utility for typical armor design problems.
INTRODUCTION
The spectrum of ballistic threats that may be encountered by ground vehicles
runs from small arms and low-velocity shrapnel from a variety of sources to high-
energy kinetic penetrators. Unfortunately, designing, integrating and fielding an
armor configuration for the highest order threat, to be used as the vehicle armor, is
clearly impractical from both weight and cost standpoints. As the US Army
acquires an expanded role in areas other than direct combat (such as Somalia or
Bosnia), armor packages more closely configured to match, rather than grossly
overmatch, the expected threat will be required. In order to meet this requirement
in a timely and affordable fashion, increased reliance is being placed on
simulation and modeling to replace the expensive process of build, shoot, build.
Ceramic Armor Materials by Design 299
To the extent authorized under the laws of the United States of America, all copyright interests in this publication are the propertyof The American Ceramic Society. Any duplication, reproduction, or republication of this publication or any part thereof, withoutthe express written consent of The American Ceramic Society or fee paid to the Copyright Clearance Center, is prohibited.
Potential armor configurations can be modeled on the computer and tested against
a large number of threats via computer simulations. In this way unacceptable or
marginally performing designs can be eliminated before committing to fabrication
and ballistic range tests, saving both time and money.
The US Army has made a major commitment to a dramatic increase in the
amount of modeling and simulation for the development of future weapon
systems. The utilization of modeling and simulation tools for end design of armor
systems is critically dependent on the accuracy of the underlying structure of such
simulations. Widespread acceptance of simulation tools hinges upon end user
trust in the predicted results. As the overall implementation of a design code can
be composed of a number of material models, it is essential that those models
accurately reflect true physical behavior. Ideally, different material models
should yield identical results, independent of the computer code used and conform
to experimental data. It is the goal of this paper to investigate the behavior of two
material models using two different computer codes.
The two ceramic models compared in this study are distinctly different: 1)
Johnson-Holmquist (JH-1) and 2) Rajendran-Grove (RG). JH-1 [1] is a
phenomenological model developed for brittle materials subjected to large strains,
high strain rates and high pressures. The equivalent strength is expressed as a
function of the pressure, strain rate, and accumulated damage; and it allows for
strength of intact and fractured material. The pressure is expressed as a function
of the volumetric strain and includes the effect of bulking for the fractured
material. JH-1 (rather than JH-2 or JH-3) was chosen for this study because it
appears to more accurately predict the SiC-B behavior. RG is a micro-crack
based constitutive model [2].
The two computer codes used in this study are the Eulerian CTH wave code
[3] and the Lagrangian EPIC hydrocode [4]. The EPIC computations were
performed with finite elements and meshless particles; the initial grids were
composed entirely of finite elements in 2D axisymmetry, and the elements were
automatically converted to particles as the elements became highly distorted [5].
The CTH computations were performed with the mix=1 option, where the yield
strength in mixed material cell is sum of volume fraction weights of individual
materials and single material cells with voids have decreased yield strength, and
the metals were modeled with Mie-Gruneisen EOS, Johnson-Cook strength and
fracture, using the same material parameters used in the EPIC simulations.
We were specifically interested in comparing computational results for two
target configurations where silicon carbide type-B ceramic is used, 1) semi-
infinite penetration in ceramics as described by Orphal and Franzen [6], and 2)
ceramic dwell as described by Lundberg, et. al [7]. These choices were made due
to the availability of the experimental data in the literature such that comparisons
could easily be made to the experimental results and because of their direct
300 Ceramic Armor Materials by Design
applicability to specific Army problems. Computations are also presented for a
tungsten penetrator impacting a steel target over a large velocity range. The
primary purpose of performing these computations was to investigate the
accuracy of the two numerical schemes using well-defined material behavior. Of
particular interest was the accuracy of the EPIC computations using the particle
algorithm since this is a relatively new technique with limited evaluation.
The following sections will present the computational results for tungsten steel
(calibration) computations; ceramic dwell (Lundberg) and ceramic penetration for
high velocities (Orphal). A brief discussion will also be presented on constant
determination.
V=1000m/sV=500m/s V=1500m/s V=2500m/sV=2000m/s V=3000m/s
P=25.2mm P=51.0mmP = 3.1 mm P=74.6mm P=78.7mmP = 67.7 mm
CTH
EPIC
P=21.0mm P=47.9mmP = 2.1mm P=73.9mm P=78.2mmP = 65.7mm
0.0
0.4
0.8
1.2
1.6
2.0
0 1000 2000 3000 4000
Imp act Velo city (m/s )
P/L
Hoh ler et al.
ep ic
cth
Tungsten alloy (D17.6)
= 17.6g/ccBHN = 406
elong(%) = 10
L = 50mm
D = 5mm
HzB,A Armor Steel
= 7.85g/ccBHN = 295
120mm
80mm
Figure 1. DOP computations in Epic and CTH
Ceramic Armor Materials by Design 301
CALIBRATION COMPUTATIONSTo first investigate the possible variations in computational results generated
by EPIC and CTH due to the numerics of the hydrocodes, computations of semi-
infinite penetration using the Johnson-Cook material model for strength and
fracture [8,9] were performed. These computations used the target geometry of
Hohler and Stilp [10]. Computations were performed over a velocity range from
500 m/s to 3000 m/s. The Brinell hardness for the targets ranged from 260-330; a
median value of 295 was used in the computations. The test configuration and
computational results are presented in Figure 1 and the Johnson-Cook constants
used in the computations are presented in Table I. The computations used the
Table I. Johnson-Cook strength and fracture constants
Mass/Thermal Properties
= 17600kg/m3
specific heat = 134.5 J/kg K
conductivity = 75.42 J/s m K
volume expansion coef. = 0.0000162
melt temperature = 1723 K
Elastic Constants
Shear Modulus (G) = 147 GPa
Shear velocity (Vs) = 2890 m/s
Bulk Modulus (K) = 287 GPa
Bulk velocity (Vb) = 4040m/s
Strength Model (Johnson-Cook)
C1 = 1.365 GPa
C2 = 0.1765 GPa
C3 = 0.016
N = 0.12
M = 1.0
Equation of State
Bulk sound velocity (Vb) = 4040m/s
Us-Up slope = 1.23
Gruneisin coefficient = 1.43
Max hydrostatic tension allowed = 68.95GPa
Fracture Model(Johnson-Cook)
D1 = 0.0
D2 = 0.33
D3 = -1.50
D4 = 0.0
D5 = 0.0
minimum fracture strain = 0.022
Spall stength = 6.757GPa
Tungsten
Mass/Thermal Properties
= 7850kg/m3
specific heat = 477.8 J/kg K
conductivity = 38.11 J/s m K
volume expansion coef. = 0.0000324
melt temperature = 1793 K
Elastic Constants
Shear Modulus (G) = 76.4 GPa
Shear velocity (Vs) = 3120 m/s
Bulk Modulus (K) = 165 GPa
Bulk velocity (Vb) = 4580m/s
Strength Model (Johnson-Cook)
C1 = 0.810 GPa
C2 = 0.5095 GPa
C3 = 0.014
N = 0.26
M = 1.03
Equation of State
Bulk sound velocity (Vb) = 4580m/s
Us-Up slope = 1.49
Gruneisin coefficient = 1.16
Max hydrostatic tension allowed = 68.95GPa
Fracture Model(Johnson-Cook)
D1 = -0.80
D2 = 2.10
D3 = -1.50
D4 = 0.002
D5 = 0.61
minimum fracture strain = 0.035
Spall stength = 5.723GPa
HzB, A Armor Steel
same material models, material input parameters and similar gridding. The depth
of penetration from the computations compared very well to the experimental
results, with the computed results predicting somewhat greater penetration (2%-
8%). However, maybe the most important result is that the CTH and EPIC
responses were very similar indicating that reasonable results can be obtained
302 Ceramic Armor Materials by Design
using very different numerics. Figure 1 also shows the final geometry of the
penetration profile for all the computations. The penetration profiles are
remarkably similar, giving additional support to expect consistent results between
the two hydrocodes.
DETERMINATION OF CERAMIC MODEL CONSTANTS
Determination of ceramic model constants for both the JH-1 and RG models is
not a straightforward process and will not be presented in detail here. The
ceramic used for this work is a hot pressed silicon carbide known as SiC-B
produced by Cercom Inc.
The process to obtain constants for the JH-1 model is presented in detail by
Holmquist [11]. Here, the majority of the constants for the JH-1 model were
measured explicitly in laboratory experiments, although two constants were
obtained by fitting model predictions to ballistic experiments. These two
constants were obtained by matching two of the dwell/penetration experiments
performed by Lundberg et. al. The process used to get these two constants was
applied in the same manner for both EPIC and CTH. The point to stress here is
that all the JH-1 constants were the same for both EPIC and CTH with the
exception of the two constants that were determined using the computations.
The constants for the RG model were obtained by matching plate impact
experiments. The same constants were used for both the EPIC and CTH
computations. The current RG model in EPIC was unable to reproduce the
Lundberg results with the RG model SiC-B constants calibrated for the low and
high velocity plate impact tests. To achieve complete dwell in the low velocity
case, the limiting crack growth rate coefficient (n1-) for mode II/III was changed
to 0.001 (a value of 0.1 was assumed to match the high velocity plate impact
data). While this change did not affect the low velocity plate impact simulation
results, the effect on the high velocity plate impact simulation results was
significant - the slower mode II/III crack growth rates resulted in delayed and
noisy spall signals that did not match the smooth spall signals measured
experimentally.
CERAMIC DWELL COMPUTATIONS
One of the most interesting ceramic characteristics is that of ceramic dwell
and interface defeat. Dwell occurs when a high velocity projectile impacts a
ceramic target and is eroded on the surface of the ceramic with no significant
penetration. If the dwell phenomenon continues until the entire penetrator is
consumed, the event is termed interface defeat (of the penetrator). Ceramic dwell
is an important characteristic of ceramic behavior and must be reproduced
computational by ceramic models. Lundberg et al. [7] demonstrated ceramic
dwell for silicon carbide (SiC-B) in a series of ballistic experiments. Three of the
Ceramic Armor Materials by Design 303
experimental results are presented in Figure 2. The two highest impact velocities
were used to get JH-1 model constants. Figure 2 shows that the JH-1 model, as
implemented in both EPIC and CTH, is capable of reproducing dwell, dwell-
penetrations transition and high velocity penetration. It should be noted that the
V = 1410m/sEPIC (JH-1) CTH (JH-1)
t = 36 s
CTH (RG)EPIC (RG)
t = 25 st = 36 st = 36 s
0
5
1 0
1 5
2 0
0 1 0 2 0 3 0 4
T im e , t ( s )
Pe
ne
tra
tio
n,
P (
m
V=2175m/s
V=1645m/s
V=1410m/s
CTH (JH-1)
0
CTH (RG)
0
5
1 0
1 5
2 0
0 1 0 2 0 3 0T im e , t ( s )
Pe
ne
tra
tio
n,
P (
4 0
m V=2175m/s
V=1645m/s
V=1410m/s
EPIC (RG)
EPIC (JH-1)
Figure 2. Comparison of experimental and computational results
two constants obtained from computations were very different for the two codes.
CTH required constants that effectively made the material softer, probably due to
the fact that the ceramic material was fixed to the confinement steel and in the
304 Ceramic Armor Materials by Design
EPIC computations it was allowed to slide. When the JH-1 values obtained using
the EPIC code were used in CTH the computations produced no penetration for
the 1645m/s experiment. Failure strain and failed yield strength values were
changed to reproduce the correct dwell phenomena. No confining pressure was
included in the simulation setup. A boundary layer around the ceramic was
attempted, but not used due to a decrease in the dwell performance. Ten cells
across the penetrator diameter were used for the mesh.
The results using the RG model also demonstrated the ability to capture the
dwell, dwell-penetration transition and the high velocity penetration, but some
modifications to the model were required. No pre-loaded confining pressure was
used in the simulations. Confining pressure was not measured in the Lundberg
experiments, but we know that some (unknown) level of confinement existed.
In the CTH simulation a single cell thickness of weak confinement material
surrounded the ceramic. This layer was identical to the actual confinement in that
it had the same EOS properties (density, pressure response, etc.), but differed in
that it had no yield strength (JC strength parameters were zero or near-zero) or
fracture strength. All simulations were axi-symmetric, utilizing square cells of
size equal to 1/8th
of the penetrator radius. As can be seen in Figure 2, the high
velocity (2175 m/s) is well represented. The duration of the transition dwell
(1645 m/s) is slightly over-predicted. The CTH-RG simulation predicted that
dwell would last for 27 s. The penetration rate after dwell is well represented.
The total dwell duration of 36 s at 1410 m/s is under-predicted, with dwell
ending at about 26 s in the simulation.
SEMI-INFINITE CERAMIC PENETRATION COMPUTATIONS
Computations were also performed into semi-infinite ceramic targets as
defined by Orphal and Franzen [6]. These computations covered a wide range of
impact velocities and were in effect “validation computations” inasmuch as the
constants were not determined from the experiments. Figure 3 presents
penetration as a function of impact velocity for the experiments and the
computations. The JH-1 model, as implemented in EPIC, produced good results
at velocities up to 2000 m/s, but tended to under-predict penetration at the high
velocities (3000 – 4000 m/s). While JH-1 in CTH produced better results at the
high velocities, it still under-predicted penetration, including a larger under-
predicted penetration at 2000 m/s. The RG model in EPIC exhibited a slightly
greater under-prediction of penetration (compared to both experimental results
and JH-1 computations). RG in CTH (still under development) predicted lower
penetration depths in the simulations of the Orphal experiments than the other
model/code implementations, but still followed the general trend of the
experimental data. Both models predicted interface defeat at 1000 m/s.
Ceramic Armor Materials by Design 305
V = 3000m/s
0.0
0.5
1.0
1.5
2.0
2.5
3.0
0 1000 2000 3000 4000 5000
Impact velocity (m/s)
P/L
Orphal and Franzen
EPIC (JH-1)
CTH (JH-1)
EPIC(RG)
CTH(RG)
Aluminum/ceramic interface
EPIC (RG) CTH (RG)CTH (JH-1)
t = 20 s
EPIC (JH-1)
t = 20 st = 20 s t = 13 s
Figure 3. Penetration computations and comparison to experiment
CONCLUSIONS
Computations were performed using the JH-1 and RG ceramic models as
implemented in the CTH and EPIC computer codes. Computations of a tungsten
rod into a steel target demonstrated that both CTH and EPIC produced very
similar results consistent with experimental data over a wide velocity range.
Computations were also performed of dwell, dwell-penetration transition and high
velocity penetration. The JH-1 model produced good results using both EPIC and
CTH. The RG model, after modifications, was able to reproduce ceramic dwell
behavior. However, in order to improve its ability to correctly predict dwell, it
has been proposed that the model should be modified to include a new "critical
shear stress" criterion that would be applied only when the ceramic material is
experiencing triaxial compressive loading (i.e., when all three principal stresses
306 Ceramic Armor Materials by Design
are compressive). Under such loading conditions, cracking could not occur unless
the maximum shear stress exceeded the critical shear stress and the Griffith
criterion was satisfied. Finally, computations were performed into semi-infinite
ceramic targets. Both models tended to under-predict the penetration into the
ceramic, but results followed the general trend of the experimental data.
Future work will include model refinement to allow better match to
experimental data and investigations considering different computational
platforms and serial versus parallel processing.
ACKNOWLEDGEMENTS
Some of this work was sponsored by the Army High Performance Computing
Research Center under the auspices of the Department of the Army, contract
number DASW01-01-C-0015. The content does not necessarily reflect the
position or the policy of the government, and no official endorsement should be
inferred. This work was supported in part by a grant of high performance
computing HPC) time from the DoD HPC Center at APG, MD.
REFERENCES
1. G. R. Johnson and T. J. Holmquist, "A Computational Constitutive Model For
Brittle Materials Subjected To Large Strains, High Strain Rates, And High
Pressures," Proceedings of EXPLOMET Conference, San Diego, (August
1990).
2. A. M. Rajendran, “Modeling the Impact Behavior of AD85 Ceramic under
Multiaxial Loading,” International Journal of lmpact Engineering, Vol. 15, pp.
749-768, (1994).
3. J. M. Mcglaun, S. L. Thompson, and M. G. Erlick, “A Three Dimensional
Shock Wave Physics Code,” International Journal of Impact Engineering,
Vo1. 10, (1990).
4. G. R. Johnson, R. A. Stryk, T. J. Holmquist and S. R. Beissel, “Numerical
Algorithms in a Lagrangian Hydrocode,” Report No. WL-TR-1997-7039
(June 1997).
5. G. R. Johnson, R. A. Stryk, L. R. Beissel, and T. J. Holmquist, “Conversion
Of Finite Elements Into Meshless Particles For Penetration Computations
Involving Ceramic Targets,” Shock Compression of Condensed Matter-2001,
in press, (2001).
6. D.L., Orphal and R.R. Franzen, “Penetration of Confined Silicon Carbide
Targets by Tungsten Long Rods at Impact Velocities from 1.5 to 4.6 km/s,”
International Journal of Impact Engineering, Vo1. 19, No. 1, pp. 1-13, (1997).
7. P. Lundberg, R. Renstrom, and B. Lundberg, “Impact of Metallic Projectiles
on Ceramic Targets: Transition Between Interface Defeat and Penetration,”
International Journal of Impact Engineering, Vo1. 24, 259-275, (2000).
Ceramic Armor Materials by Design 307
8. G. R. Johnson and W. H. Cook, “A Constitutive Model and Data for Metals
Subjected to Large Strains, High Strain Rates, and High Temperatures,”
Proceedings of Seventh International Symposium on Ballistics. The Hague,
The Netherlands, (April 1993).
9. G. R. Johnson and W. H. Cook, “Fracture Characteristics of Three Metals
Subjected to Various Strains, Strain Rates, Temperatures, and Pressures,”
Engineering Fracture Mechanics, Volume 21, (1985).
10. C. E. Anderson, Jr., B. L. Morris and D. L. Littlefield, “A Penetration
Mechanics Database,” SwRI Report 3593/001, (January 1992).
11. T. J. Holmquist and G.R. Johnson, “Response of Silicon Carbide to High
Velocity Impact,” submitted for publication, Journal of Applied Physics,
(2001).
308 Ceramic Armor Materials by Design
MODELING CERAMIC DWELL AND INTERFACE DEFEAT
Timothy J. Holmquist and Gordon R. Johnson
Network CS/Army High Performance Computing Research Center
1200 Washington Avenue South
Minneapolis, MN 55415
ABSTRACT
This paper presents computational modeling of ceramic dwell and interface
defeat, using the EPIC code and the JH-1 constitutive model for ceramics.
Computations are presented for various projectiles impacting various silicon
carbide (SiC-B) targets. The computational results are shown to provide good
agreement with experimental data in the literature. Also included are the JH-1
constants for SiC-B, the procedure used to determine the constants, and a
description of some important computational features involving finite elements
and meshless particles.
INTRODUCTION
Ceramic materials are generally strong in compression, weak in tension, and
can have considerable strength after failure when they are under compression.
They have been used as armor materials for many years. More recently,
experimental data have been presented by Lundberg et al. [1] that show how
silicon carbide targets can be configured to defeat tungsten and molybdenum rods
at significant impact velocities. Other researchers have demonstrated this same
effect for other ceramics, but this paper will focus only on the silicon carbide
targets.
JH-1 CERAMIC MODEL AND CONSTANTS
The JH-1 constitutive model for ceramics, and the associated constants for
SiC-B, are shown in Figure 1. The model consists of an intact strength and a
failed strength that are functions of the pressure, the strain rate, and the damage.
Pressure, bulking and damage are other aspects of the model. This is the first of
two closely related models, JH-1 [2] and JH-2 [3], presented by Johnson and
Holmquist. One of the primary differences between the two models is that the
JH-2 model allows the strength to degrade gradually as the damage is
Ceramic Armor Materials by Design 309
To the extent authorized under the laws of the United States of America, all copyright interests in this publication are the propertyof The American Ceramic Society. Any duplication, reproduction, or republication of this publication or any part thereof, withoutthe express written consent of The American Ceramic Society or fee paid to the Copyright Clearance Center, is prohibited.
accumulated, rather than soften/fail instantaneously after it is fully damaged, as is
done in the JH-1 model. For SiC-B the JH-1 model appears to be better suited to
represent the strength and interface defeat characteristics of SiC-B. Apparently
the JH-2 approach, with gradual softening, does not provide the constant target
resistance required for dwell and interface defeat.
Fai
lure
str
ain,
pf
T P3
D= p/ pf
Pressure, P Volumetric strain,
Pre
ssure
, P
D<1.0
D=1.0
P
T
Density o = 3215kg/m3
Shear modulus G = 193GPa
Tensile strength T = 0.75GPa
Intact strength S1 = 7.1GPa
Intact strength P1 = 2.5GPa
Intact strength S2 = 12.2GPa
Intact strength P2 = 10.0GPa
Strain rate C = 0.009
Failed strength Sfmax= 1.3GPa
Failed strength = 0.40
Bulk modulus K1 = 220GPa
Pressure K2 = 361GPa
Pressure K3 = 0GPa
Bulking factor = 1.0
Damage = 0.012Eq
uiv
alen
t S
tres
s,
T P1 P2
S2
S1
Pressure, P
*>1.0.
*=1.0.
Intact Material (D<1.0)
Failed Material (D=1.0)
P=K1 +K22+K3
3
fmaxS
f
max
)TP/( 3
f
max
Figure 1. JH-1 model and constants for silicon carbide.
The intact strength and compressibility constants were obtained from the plate
impact data of Feng et al. [4], Grady and Moody [5], Dandekar and Bartkowski
[6], and the uniaxial compression data of Pickup and Barker [7]. A very
important characteristic of the Feng et al. data is that both longitudinal and lateral
stresses are provided, and this allows the complete stress state (both strength and
pressure) to be determined. The slope constant for the failed material, , is
provided by Klopp and Shockey [8].
310 Ceramic Armor Materials by Design
The maximum strength of the failed material, cannot be obtained
directly due to the lack of experimental data, but must instead be inferred by
varying until computed penetration computations match the results of
experimental data. The damage constant,
,max
fS
fSmax
, is obtained in a similar manner. The
two experimental data points used to obtain these constants are the penetration of
a tungsten rod into a silicon carbide target at 2175 m/s, and the duration of the
dwell of a tungsten rod onto a ceramic target at 1645 m/s [1]. Additional details
concerning the model and the constants are provided elsewhere [9].
EXAMPLES
The following examples all use the silicon carbide constants for the JH-1
model as shown in Figure 1. The computations were performed with finite
elements and meshless particles; the initial grids were composed entirely of finite
elements in 2D axisymmetry, and the elements were automatically converted to
particles as the elements became highly distorted [10].
Figure 2 shows computational results compared to experimental results
provided by Lundberg et al. [1]. The tests were performed by impacting a
confined ceramic target onto a long stationary rod (of tungsten or molybdenum).
Only three test results are shown for each of the two rod materials, although
additional test results are reported by Lundberg et al. These three results
correspond to interface defeat (where the rod does not penetrate the ceramic),
dwell and penetration (where the rod dwells on the surface of the ceramic before
it begins to penetrate), and penetration (where the rod penetrates with minimum
dwell).
For the tungsten tests, the two higher impact velocities (1645 m/s and 2175
m/s) were used to determine the strength of the failed material, , and the
damage constant,
fSmax
. Although there is some coupling between these two
constants, the penetration rate/depth is most dependent on and the dwell is
most dependent on
fSmax
. It can be seen there is very good agreement between the
computed results and the experimental results for interface defeat, dwell, and
penetration. Although the tungsten experimental results were used to determine
some of the constants, the molybdenum experimental results were not used. It
should be noted that the molybdenum test data for interface defeat were for an
impact velocity of 2030 m/s, whereas the corresponding computed results are for
a slightly reduced impact velocity of 2000 m/s.
Figure 3 shows computed results for the interface defeat of the tungsten rod for
an impact velocity of 1410 m/s. For the three geometry plots in the upper portion
of the figure, the tungsten rod is represented by the darkened elements and
particles; and the steel plug, steel tube and ceramic are represented by various
Ceramic Armor Materials by Design 311
colors of elements and particles. The black lines define the outlines between the
materials, and in other instances they represent an interface between the elements
and particles. For this case the ceramic remains intact, while the defeated rod
moves radially outward along the top surface of the ceramic until it is contained
by the steel case. The distribution of damage is shown in the lower right portion
of the figure, where it can be seen that most of the ceramic under the rod is only
partially damaged, and this enables the ceramic to remain intact and to defeat the
rod.
Steel plug
Silicon carbide
L = 20mm
D = 20mm
Steel plug
Hig
h s
tren
gth
ste
el t
ub
e
Tungsten or
Molybdenum
rod
L = 80mm
D = 0.5mm
0
5
10
15
20
0 10 20 30 40
0
5
10
15
20
Lundberg et al.
Tungsten rod
Molybdenum rod
V=2175m/s
V=1645m/s
V=1410m/s
V=2535m/s
V=2090m/s
V=2000m/s
Pen
etra
tion
in
to c
eram
ic(m
m)
(a) (b)
00
Computed
results
Time, t (microsecond)
Figure 2. Comparison of computational results and experimental results for tungsten
and molybdenum rods impacting a confined silicon carbide target at various velocities.
(a) Initial 2D geometry and (b) comparison of computed and experimental results.
Figure 4 shows the computed response for a slightly higher impact velocity of
1645 m/s. Here the tungsten rod dwells for a short time and then penetrates the
ceramic. When the ceramic material directly under the impacting rod becomes
completely damaged and fails, the dwell ceases and the penetration begins. The
distribution of damage for the transition between dwell and penetration is shown
in the lower right portion of the figure.
A comment should be made concerning an important advantage of converting
distorted elements into particles rather than simply eroding (or removing) the
distorted elements. When an element is eroded it introduces a void which allows
312 Ceramic Armor Materials by Design
surrounding material to expand into the void and to lose pressure as it expands. If
the material strength or failure characteristics are pressure dependent (as they are
for ceramics) then the pressure drop can lead to lower strength and/or increased
damage.
Figure 3. Computed results for a tungsten rod impacting a confined silicon carbide target
at 1410m/s. (a) penetration of steel cover at 12 s, (b) dwell at 20 s, (c) dwell at 36 s,
(d) close up of material flow at 36 s, and (e) close up of material damage at 36 s.
The final example involving dwell is shown in Figure 5. Here a pointed steel
projectile impacting a thin plate composed of a silicon carbide layer (6.35 mm
thick) over a 6061-T6 aluminum layer (6.35 mm thick) is investigated. An impact
velocity of 700 m/s is shown, as it represents the computational ballistic limit.
This is only slightly higher than the experimental ballistic limit of 660 m/s
reported by Wilkins [11]. It can be seen that there is significant dwell during the
initial 10 to 20 and that the damage pattern in the ceramic forms in a conicals
Ceramic Armor Materials by Design 313
pattern. The silicon carbide tested by Wilkins was not the SiC-B for which the
constants were obtained, and it is not known if the experimental results would be
significantly different for SiC-B. Nevertheless, the basic trends and mechanisms
appear to be well represented by the computations.
Figure 4. Computed results for a tungsten rod impacting a confined silicon carbide
target at 1645m/s. (a) dwell at 12 s, (b) transition from dwell to ceramic penetration at
18 s, (c) ceramic penetration at 30 s, (d) close up of material flow at 18 s, and (e)
close up of material damage at 18 s.
SUMMARY AND CONCLUSIONS
This paper has demonstrated the computational capability to simulate ceramic
dwell and interface defeat. The specific form of silicon carbide, known as SiC-B,
has been characterized from experimental data in the literature and put into the
form of the JH-1 constitutive model for ceramics. Many of the constants were
determined explicitly from the experimental data, but some of the constants for
314 Ceramic Armor Materials by Design
damaged/failed material were inferred from ballistic penetration data. There was
generally very good agreement between the computational results and the
experimental results, even for the problems that were not used to determine the
constants.
Figure 5. Computed results for a steel projectile impacting a thin, layered target of
silicon carbide and aluminum at 700 m/s. (a) Material flow at 10 s, 20 s and 200 s
after projectile impact, and (b) material damage at 10 s, 20 s and 200 s after
projectile impact.
ACKNOWLEDGEMENTS
This work was sponsored by the Army High Performance Computing
Research Center under the auspices of the Department of the Army, contract
number DASW01-01-C-0015. The content does not necessarily reflect the
Ceramic Armor Materials by Design 315
position or the policy of the government, and no official endorsement should be
inferred. The authors would also like to thank D. W. Templeton (Army Tank
Automotive Research, Development, and Engineering Center) for his
contributions to this work.
REFERENCES1 P. Lundberg, R. Renstrom, and B. Lundberg, “Impact of metallic projectiles
on ceramic targets: transition between interface defeat and penetration,”
International Journal of Impact Engineering, 24, p. 259, (2000). 2 G. R. Johnson and T. J. Holmquist, “A computational constitutive model for
brittle materials subjected to large strains, high strain rates, and high pressures,”
Proceedings of EXPLOMET Conference, San Diego, August 1990. 3 G. R. Johnson and T. J. Holmquist, “An improved computational constitutive
model for brittle materials,” High Pressure Science and Technology – 1993,
Edited by S. C. Schmidt, J. W. Schaner, G. A. Samara, and M. Ross, AIP 1994.
R. Feng, G. F. Raiser, and Y. M. Gupta, “Material strength and inelastic
deformation of silicon carbide under shock wave compression,” Journal of
Applied Physics, 83, p.79, (January 1998).
4
D. E. Grady and R. L. Moody, “Shock compression profiles in ceramics,”
Report No. SAND96-0551, Sandia National Laboratory, March 1996.
5
D. P. Dandekar and P. T. Bartkowski, “Tensile strengths of silicon carbide
(SiC) under shock loading,” Report No. ARL-TR-2430, Army Research
Laboratory, March 2001.
6
I. M. Pickup and A. K. Barker, “Deviatoric strength of silicon carbide subject
to shock,” Shock Compression of Condensed Matter – 1999, Edited by M. D.
Furnish, L. C. Chhabildas, and R. X. Hixon, p. 573, (AIP, 1999).
7
8 R. W. Klopp and D. A. Shockey, “The strength behavior of granulated silicon
carbide at high strains rates and confining pressure,” Journal of Applied Physics,
70, p.7318, (December 1991).
T. J. Holmquist and G. R. Johnson, “Response of silicon carbide to high
velocity impact,” Submitted for publication.
9
G. R. Johnson, R. A. Stryk, S. R. Beissel, and T. J. Holmquist, “Conversion
of finite elements into meshless particles for penetration computations involving
ceramic targets,”Shock Compression of Condensed Matter – 2001, in press.
10
M. L. Wilkins, “Fourth progress report of light armor program,” Report No.
UCRL-50694, Lawrence Radiation Laboratory, 1969.
11
316 Ceramic Armor Materials by Design
3D FINITE ELEMENT ANALYSIS OF IMPACT DAMAGE IN METALLIC
AND CERAMIC TARGETS
Fenghua Zhou and Jean-Francois Molinari
Department of Mechanical Engineering, Johns Hopkins University
232 Latrobe Hall, 3400 N. Charles St., Baltimore, Maryland, 21218
ABSTRACT
A 3D explicit FEM package has been developed to analyze the deformation
and failure process of materials under impact loading. The package is constituted
of a finite deformation plastic model, a frictional contact algorithm, an adaptive
meshing capability, and a module of dynamically inserted cohesive elements to
simulate dynamic crack propagation. In this paper, we present three simulations
that investigate dynamic crack propagation in a metallic and a ceramic material.
The quantitative and qualitative agreements between experimental and simulation
results are discussed.
INTRODUCTION
The investigation of the dynamic response and the failure process of metallic
and ceramic targets under impact loading is a crucial issue to the designing of
armor and anti-armor systems. Numerical techniques are broadly applied in this
investigation to analyze and to simulate expensive experimental results. Although
many large-scale, general-purpose codes for conducting structural impact analysis
have been developed, the problem of modeling material deformation/failure
processes under intensified loading is still a challenging task. These processes are
accompanied with the mechanisms of large-deformation high strain-rate
plasticity, heat generation and conduction, and dynamic fracture and
fragmentation. Moreover, the need of using lighter and stronger composite
materials has lead to the development of complex materials such as metal-matrix
or ceramic-matrix composite materials. The combination of a variety of physical
phenomena with a range of microstructures renders the interpretation of impact
mechanisms difficult. As a design tool and as a methodology to reach a physical
understanding, a multi-physics finite element tool is of great importance.
We have developed a three-dimensional explicit finite element analysis
package. The package includes finite deformation plasticity, frictional contact,
Ceramic Armor Materials by Design 317
To the extent authorized under the laws of the United States of America, all copyright interests in this publication are the propertyof The American Ceramic Society. Any duplication, reproduction, or republication of this publication or any part thereof, withoutthe express written consent of The American Ceramic Society or fee paid to the Copyright Clearance Center, is prohibited.
heat generation and conduction, and adaptive meshing [1, 2]. Molinari et al. [1]
studied the impact erosion in mild-steel targets by hard-steel projectiles. The
range of impact velocity varied from 200m/s to 2000m/s and the impact angle
from glancing to normal penetration. The calculations highlighted that the friction
resistance increases with the sliding velocity for the low range of impact speeds.
Another outcome of the calculations is that at 45 degrees impact frictional heating
becomes dominant. However, crack propagation was not modeled in [1]. In this
paper, we highlight the addition of cohesive elements to the existing capability.
Cohesive zone models can explicitly describe the crack initiation and propagation
process. Several simulations of the dynamic fracture phenomena were conducted
to verify and validate the methodology. We begin with a brief description of the
relevant aspects of the material and numerical model. Then, to illustrate the
methodology, we simulate the dynamic bursting of a ceramic ring under
centrifugal force. Impact three-point-bending simulations performed on metallic
and ceramic materials constitute the core of the paper.
FINITE ELEMENT ANALYSIS AND COHESIVE MODELS
In our analysis, the volume of the structure is meshed by 10-node (quadratic)
tetrahedral elements. The simulation of the dynamic process is conducted by
using the second order accurate explicit form of Newmark’s algorithm ([Hughes,
[3]):
)(2
1)(
2
1
11
int
11
1
1
2
1
nnnn
n
ext
nn
nnnn
t
tt
aavv
FFMa
avdd
(1)
where the subscript n denotes variables at the nth
time step; t is a fraction of the
critical time step; d, v and a are nodal displacement, velocity and acceleration
vectors; M is the lumped mass matrix, Fext
and Fint
are the external and internal
nodal forces.
We use an isotropic J2 flow material model that includes thermal softening,
power-law strain hardening and strain-rate hardening (Cuitino et al., [4]):
n
p
p
refmelt
ref
y
m
pp
p
TT
TTg
Tg1
0
0
11
1),(
(2)
where is the effective Mises stress, the effective plastic strain, the
effective plastic strain rate, g the flow stress and T the temperature. The material
p p
318 Ceramic Armor Materials by Design
parameters include the initial yield stress, the reference plastic strain,
the reference plastic strain rate, m the rate sensitivity exponent, n the hardening
exponent, T the reference temperature, T the melting temperature and the
thermal softening exponent.
y
2
p
0
Rigid
Smith
p
0
ref melt
g:d >0
The key issue in this paper is the simulation of dynamic crack propagation.
We follow a cohesive element type approach. The cohesive zone concept was
firstly introduced by Dugdale [5] and Barrenblatt [6], who independently assumed
that a cohesive region exist at the crack tip, where stress singularity is eliminated
by the distribution of the cohesive forces. The cohesive zone concept was
implemented into numerical analysis to explicitly simulate crack propagation. The
work of Xu and Needleman [7, 8], Camacho and Ortiz [9, 10] demonstrated
successful use of the cohesive element in 2D cases. The work of Pandolfi et. al.
[11, 12], and Ruiz et. al. [13, 14] successfully implemented cohesive elements
into 3D analysis in a range of applications. Other recent relevant numerical
investigations include the effect of microstructure on the dynamic failure process
[15, 16].
In our calculations we insert 12-nodes triangular cohesive elements between
two neighboring tetrahedral elements, Fig. 1a.
Fig.1 (a) Cohesive element between two tetrahedral elements;
(b) Two irreversible cohesive models: Smith-Ferrante law and Rigid-linear law
c
c
c
c
O penin Closing : d 0
The tractions (cohesive force) between the glued faces (S+ and S- in Fig. 1a)
are functions of their relative distance. These functions are called cohesive laws,
and they express the energy and the forces needed to open the cohesive element.
Two cohesive laws frequently used are shown in Fig. 1b. They are the irreversible
exponential decaying and the irreversible linear decaying functions [11, 12].
Irreversibility signifies that the damage in a given cohesive element cannot be
recovered. The area under the curves of Fig. 1b is the fracture energy, which is
needed to fully open a unit area of crack surface.
In the models sketched in Fig. 1b the fracture energy takes a simple form:
modellinear-
modelFerrante-
5.0cc
cc
cc
eG (3)
Ceramic Armor Materials by Design 319
where Gc is the fracture energy, c is the surface energy, c is the critical opening
stress, and c is the fracture strength. Both models have their merits and demerits.
The Smith-Ferrante law is physically reasonable but inserting this element to the
structure may modify structural compliance; the rigid-linear law does not modify
the structural elastic properties, but it needs to be introduced dynamically. Its use
necessitates some computational effort. In our calculations, the simulation results
are not significantly affected by the choice of the cohesive law.
For conciseness, the remaining components of the methodology, adaptive
meshing, heat generation and conduction and frictional contact, are not described
here. A detailed description is contained in [1, 2].
DYNAMIC FRACTURE OF A ROTATING CERAMIC DISK
As an illustration of the use of cohesive
elements, we run a test simulation of the
burst of a ceramic ring under centrifugal
forces. The rotating ring experiments are a
standard test to derive mechanical
properties of materials [17]. Our model
consists of 5438 nodes and 2614 tetrahedral
elements, Fig. 2. The cohesive elements are
inserted when a critical centrifugal load is
reached. The disk is constituted of Si3N4,
whose properties are listed in Table. I. Note
that plasticity was neglected as the material
considered is brittle. The data was obtained from NIST
[http://www.ceramics.nist.gov/srd/summary/ftgsin.htm]. Two numerical tests,
with different fracture properties, Type-1 and Type-3 (Table. I), are conducted.
They reflect the variations in materials data handbooks and assess the numerical
sensitivity to materials data. In both types we keep the critical tensile stress ( c) at
a constant value of 450 MPa. However, the fracture energy (Gc) varies from 100
N/m for Type-1 to 200 N/m for Type-3.
Table I. Mechanical properties of Si3N4
Fracture Properties
(Irreversible Linear Decreasing Cohesive Law)Densi
ty
Elastic
Type-1 Type-2 Type-3 Type-4
Gc = 200 N/m Gc = 100 N/m
kg/m3
E
GPa c
Gpac
mc
GPac
mc
GPac
mc
GPac
m
3300 320 0.23 0.45 0.889 1.0 0.4 0.45 0.444 1.0 0.2
Fig. 2 Ceramic ring before burst
320 Ceramic Armor Materials by Design
As the fracture stress is kept to a constant value in both tests, the rotating speeds
and the time at which the ceramic ring bursts are identical (about 47500rpm). Fig.
3 shows the initiation and crack propagation of a Type-1 ring. Clearly, multiple
cracks initiate at the inner rim, where the hoop stress is maximum. The cracks
then expand outwards and in some instances branching occurs when approaching
the outer rim. The final shapes and sizes of the fragments are shown in Fig. 4. The
two tests highlight that the fracture energy directly affects the number of
fragments. The number of cracks and fragments increases with a decreasing
fracture energy. This is qualitatively in agreement with experimental observations
[17].
DYNAMIC 3 POINT BENDING FRACTURE TEST ON SIC/AL MATERIAL
We now focus on 3-point bending
calculations in which ductile and brittle
materials are impacted by a Kolsky bar. In
this section, we simulate the dynamic
fracture process of silicon carbide particle
reinforced aluminum alloy (SiCp/Al). The
dynamic mechanical properties of this
material have been thoroughly investigated
by Li et. al. [18, 19]. The dynamic fracture
behavior of the material was also
experimentally studied using a Kolsky-bar
testing system [20]. The experimental setup is shown in Fig. 5: a 3-point bending
specimen is sandwiched between an input bar and an output tube. When a
projectile impacts the input bar, a compressive stress wave propagates along the
bar. This wave is transmitted to the specimen. The loading history on the
impacted point of the specimen can be measured by using strain-gages adhered to
the bar and tube. During the experiments, the local strains near the crack tip, and
the crack tip opening distance (CTOD) are monitored. Four tests were performed,
#6, #7, #8 and #9, where the impact speeds of the projectile were respectively
6m/s, 16.8m/s, 16.9m/s and 8.5m/s.
Fig. 3 The propagation of cracks in a
ceramic ring (Type-1 material) Fig. 4 Ceramic rings after burst (50 s):
(a) Gc=200N/m; (b) Gc=100N/m
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������������
C rack-Tip Strain G age
C TO D M arking Points
Incident B arProjectile O utput Tube
v0 Strain G age
Fig.5 Kolsky-bar system for 3-
point bending fracture tests
Ceramic Armor Materials by Design 321
The FEM model is shown in Fig. 6a. The specimen is supported on the output
tube, which is modeled as fixed in the y-direction. The forces applied on the
specimen, (t) are dependent on the structural response of the specimen.
Therefore at the impact zone we apply a boundary condition of the form:
)()(2)( tcvtt inc (4)
where inc(t) is the input stress wave (experimental data), v(t) is the velocity of
loading point (computed), and c are the density and elastic wave speed of the
input bar. In our analysis, the incident wave is used as the input data. Eqn (4) is
introduced into the explicit algorithm to calculate the response of the specimen.
The reflective wave is an output of the calculations. This wave is compared to the
experimental data as a validation of the numerical results.
The mesh of the specimen is shown in Fig. 6b. It contains 24609 nodes and
16251 tetrahedral elements. A simulation using a finer mesh (54325 nodes, 37010
elements) was conducted. The numerical results were similar to the one obtained
with the coarser mesh. The material data used, is gathered in Table II ([18]). The
temperature is taken to be equal to Tref so that thermal softening is not considered
in our analysis.
Table II. Mechanical Properties of SiCp/Al
Den-
sity
Elastic Plastic Strain
Hardening
Strain Rate
Hardening
Fracture Properties
(Irreversible Linear)
kg/m3
E
GPa y
MPa0p n d 0
p/dt
1/s
m Gc
N/mc
GPa c
m
2738 102 0.29 210 1.556E-2 3.76 1.466E5 2.22 2306 1.02 4.52
X Y
Z
Fig.6 (a) Model for 3-point bending dynamic fracture tests; (b) Mesh
322 Ceramic Armor Materials by Design
The results of the calculated and
experimental reflection waves are compared
in Fig. 7. The quantitative and qualitative
agreement is good. It is noteworthy that the
material parameters were not fitted to match
the experimental structural response.
The computed local strains, which
contain an elastic and a plastic part are
compared to the experimental results in Fig.
8. Note that in each test, the strain increases
to a maximum, at which a crack propagates,
and subsequently decreases to reach a
constant value (the irreversible plastic part).
The numerical results compare well with experiments in two aspects. First, the
times at which the crack-tip strains drop, match the experimental data, which
implies that the crack initiation time is accurately simulated. Second, the
magnitudes of the strain drop after crack propagation quantitatively agree with the
experimental data (Fig. 8a). However, the experimental peak strains are about 10-3
lower than the simulation results. The existence of residual strains at the
specimen’s crack tip may be the reason of such discrepancy. The crack tip
opening distances are compared in Fig. 8b. A quantitative agreement is observed.
An example of crack propagation is shown in Fig. 9. It can be seen that the
crack front is curved because of 3D effects. The average velocity of the main
crack propagation is shown in Fig. 10. The propagation velocity of the main crack
increases with the increase of impact velocity.
0.00 0.02 0.04 0.06 0.08 0.10 0.12 0.14
0
2
4
6
8
Velo
cit
y (
m/s
)
Time (ms)
#6
#9
#7
#8
Points: Exp. Data
V_Ref #6
V_Ref #7
V_Ref #8
V_Ref #9
Fig. 7 The reflection waves from
simulations and experiments
0.00 0.02 0.04 0.06 0.08 0.10
0.0000
0.0005
0.0010
0.0015
0.0020
0.0025
0.0030
Points: Exp. Data
Str
ain
Near
Cra
ck T
ip
Time (ms)
Strain #6
Strain #7
Strain #8
Strain #9
0.00 0.01 0.02 0.03 0.04 0.05 0.06 0.07 0.08 0.09 0.10
0.00
0.01
0.02
0.03
0.04
0.05
0.06
0.07
0.08
0.09
0.10
Points: Experimental Data
CT
OD
(m
m)
Time (ms)
CTOD #6
CTOD #7
CTOD #8
CTOD #9
Fig. 8 Comparisons of the local variables from simulations and the experiments
(a) The crack-tip strains; (b) The crack tip opening distances
Ceramic Armor Materials by Design 323
It is noteworthy that the cohesive laws incorporated in the present analysis are
rate-independent. Nevertheless, the phenomenological rate-dependent dynamic
fracture process can be simulated. The reason is that cohesive fracture models
have an intrinsic time scale, which is linked to the characteristic length scale ( c)
and the wave velocity (c), Camacho et al. [9]. This property is demonstrated in
our simulations. As shown in Fig. 7 and 8, the simulations using the same set of
material data match experiments carried out at different loading rates.
DYNAMIC FRACTURE OF CERAMIC MATERIAL, VIRTUAL TEST
Having validated the implementation of cohesive elements, we now simulate
a dynamic fracture test on ceramics. No further comparison with experiments is
carried out. Thus, in essence, our following simulations constitute a virtual test,
which may be used to scale and understand experiments as well as evaluate
dynamic material properties. The specimen material is changed to silicon nitride.
We consider the material to be elastic-brittle under loading. Since an accurate
dataset of micro-cracking parameters is not available, four material types are
assumed, as listed in Table I. The specimen is loaded by the Kolsky-bar with an
incident stress wave equal to the one in Test #6 of the previous section. The
response of the specimen, and the virtual experimental measurement, are
predicted as following.
Fig. 11 shows the
reflection wave and the
loading history on the
ceramic specimen. Only
minor differences are
seen for the four types
of materials. This is
reasonable since the
stress wave
measurement contains
Fig. 9 Crack propagation in SiCp/Al specimen
(Test #8)
6 8 10 12 14 16 18
120
130
140
150
160
170
180
190
200
210
Avera
ge C
rack V
elo
cit
y (
m/s
)
Kolsky-bar Impact Velocity (m/s)
Crack Velocity
Fig. 10 Crack velocity in
SiCp/Al specimen
0.00 0.01 0.02 0.03 0.04 0.05
-10
-8
-6
-4
-2
0
2
4
6
8
Kolsky Bar Measurement
Velo
cit
y (
m/s
)
Time (ms)
Incident
Type-1 reflecion
Type-2 reflection
Type-3 reflection
Type-4 reflection
0.00 0.01 0.02 0.03
0
200
400
600
800
1000
1200
1400
1600
Lo
ad
(M
Pa)
Time (ms)
Type-1 Load
Type-2 Load
Type-3 Load
Type-4 Load
Fig. 11 Kolsky-bar recordings: (a) Incident and
reflection waveform; (b) Loading history
324 Ceramic Armor Materials by Design
the response of the whole specimen structure.
On the other hand,
the recordings of crack-
tip strains and CTOD,
shown in Fig. 12,
demonstrate significant
differences for the four
types of material
properties. The reason is
that such recordings are
local to the crack tip
zone and therefore are more sensitive to the initiation of crack propagation. The
initiation time of the crack is strongly affected by the microscopic fracture
properties: the larger the fracture energy (Gc) or the larger the critical strength
( c), the later the crack initiates.
An example of crack propagation is shown in Fig. 13. It is seen that the front
of the crack is straighter in comparison to the case of SiCp/Al material (Fig. 9).
The reason is that the brittleness of the material diminishes the 3D effects.
The locations of the
crack front, and the
average crack velocity
are shown in Fig. 14.
The speed of the crack
propagation in ceramics
is about 1000m/s,
which is much higher
than for the SiCp/Al
specimen. The velocity
of the crack not only depends on the fracture energy, but also depends on the
loading history of the specimen. An accelerating-decelerating crack behavior can
be seen from the figure, as the external load on the specimen drops (see Fig. 10b).
We also simulate a case where the fracture energy is not constant, but decreases
linearly along the specimen’s width. The results are also shown in Fig. 14. In this
0.00 0.01 0.02
0.0000
0.0001
0.0002
0.0003
Cra
ck T
ip S
train
Time (ms)
Type-1 Strain
Type-2 Strain
Type-3 Strain
Type-4 Strain
0.000 0.005 0.010 0.015 0.020 0.025
0.000
0.005
0.010
0.015
0.020
Cra
ck T
ip O
pen
ing
Dis
tan
ce (
mm
)
Time (ms)
Type-1 CTOD
Type-2 CTOD
Type-3 CTOD
Type-4 CTOD
Fig. 12 Local quantity recordings: (a) Crack-tip strains;
(b) Crack-tip opening distances
Fig. 13 Crack propagation in ceramic specimen (Type-1 material)
Fig. 14 Crack front locations and crack velocities
(Type-1 and Type-3 material)
Ceramic Armor Materials by Design 325
case, the velocity of the crack becomes relatively uniform. We will address the
issue of a limiting crack velocity and its practical consequences on design of
crack-resistant armors in future calculations.
CONCLUDING REMARKS
We have developed an explicit dynamic element package, which includes
large deformation plasticity, contact, adaptive meshing and dynamic insertion of
cohesive elements. This paper highlights the introduction of 3D cohesive
elements. We simulate three types of dynamic fracture phenomena: the
fragmentation of a brittle ceramic ring, crack propagation in a ductile metallic
specimen and crack propagation in a brittle ceramic specimen. The first
simulation illustrates how cohesive elements can be used to model explicitly
crack propagation and the complex associated fragmentation process. In the
second simulation, the results are compared to the experiments and quantitative
agreements are obtained. In this light, it validates the methodology. The intrinsic
time scale of the cohesive elements, which permits the reproduction of
experiments at various loading rates, is highlighted. The last simulation
constitutes a virtual test capability, which can be used to evaluate material
properties and design structures in which the crack velocity needs to be
controlled.
ACKNOWLEDGEMENT
The research is sponsored by Army Research Lab under contract
DAAD19012003. The authors would like to thank Professor K.T. Ramesh and Dr.
Y. Li of the Johns Hopkins University for the invaluable discussions.
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D.S. Dugdale, “Yielding of Steel Sheets Containing Slits”, J. Mech. Phys.
Solids, 8, 100-104 (1960).
326 Ceramic Armor Materials by Design
6G.I. Barrenblatt, “The Mathematical Theory of Equilibrium of cracks in
Brittle Fracture”, Adv. Apply. Mech., 7, 55-129 (1962). 7
X.-P. Xu and A. Needleman, “Numerical Simulations of Fast Crack Growth
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X.-P. Xu and A. Needleman, “Numerical Simulations of Dynamic Crack
Growth Along an Interface”, Int. J. Fracture, 74, 289-324 (1996). 9
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in Brittle Materials”, Int. J. Solids Structures, 33, 2899-2938 (1996). 10
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Penetration of Metallic Targets”, Comput. Meth. Appl. Mech. Engng, 142, 269-
301 (1997). 11
A. Pandolfi, P. Krysl and M. Ortiz, “Finite Element Simulation of Ring
Expansion and Fragmentation: The Capturing of Length and Time Scales
Through Cohesive Models of Fracture”, Int. J. Fracture, 95, 279-297 (1999). 12
A. Pandolfi, P.R. Guduru, M. Ortiz and A.J. Rosakis, “Three Dimensional
Cohesive-Elements of Dynamic Fracture in C300 Steel”, Int. J. Solids Structures,
37, 3733-3760 (2000). 13
G. Ruiz, M. Ortiz and A. Pandolfi, “Three Dimensional Finite-Element
Simulation of the Dynamic Brazilian Tests on Concrete Cylinders”, Int. J. Numer.
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G. Ruiz, M. Ortiz and A. Pandolfi, “Three Dimensional Cohesive Modeling
of Dynamic Mixed-Mode Fracture”, Int. J. Numer. Meth. Engng, 52, 97-120
(2001).15
J. Zhai and M. Zhou, “Finite Element Analysis of Micromechanical Failure
Modes in a Heterogeneous Ceramic Material System”, Int. J. Fracture, 101, 161-
180 (2000). 16
P.D. Zavattieri and H.D. Espinosa, “Grain Level Analysis of Crack
Initiation And Propagation in Brittle Materials”, Acta. Mater. 49, 4291-4311
(2001).17
R. Hashimoto, A. Ogawa, T. Morimoto and M. Yonaiyama, “Spin Tests of
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Y. Li, K.T. Ramesh and E.S.C. Chin, “The Compressive Viscoplastic
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Alloy Matrix”, Int. J. Solids Structures, 37, 7547-7562 (2000) 19
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Y. Li, K.T. Ramesh and E.S.C. Chin, “A Simple Approach to the
Measurement of Dynamic Fracture Toughness”, to be published
Ceramic Armor Materials by Design 327
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A NUMERICAL INVESTIGATION OF MICROCRACKING DIFFUSION IN SANDWICHED GLASS PLATES
Z. Chen and L. Shen G.I. Kanel and S.V. Razorenov Dept. of Civil and Environmental Engr. Inst. for Chemical Physics Research University of Missouri-Columbia Russian Academy of Sciences Columbia, MO 65211-2200, USA Moscow Region, Russia
ABSTRACT Based on the previous research on modeling and simulation of the failure wave phenomenon as observed in shocked glasses, a numerical investigation is conducted here to simulate microcracking diffusion in sandwiched glass plates. The essential assumptions made are that the deviatoric elastic strain energy in the intact material is converted into the volumetric potential energy in the comminuted and dilated material during the time- and space-dependent microcracking diffusion process, and that each surface of the glass plates introduces additional microcracking sources. The simulation results appear to match the available experimental data very well in the loading phase. Future work is discussed based on the current results.
INTRODUCTION
Under plate impact, some brittle solids may undergo elastic deformations at the shock wave front, and fail catastrophically at a distinctly later time if the shock stress is near but below the apparent Hugoniot elastic limit (HEL). The phenomenon has therefore been interpreted as the result of a slowly propagating failure wave in the shocked solids. Since Brar et al.1 and Kanel et al.2 reported the formation and propagation of failure waves in shocked glasses, continued efforts have been made to explore this interesting physical phenomenon.3-16 However, no consensus can be made at the moment on the exact physics behind this failure wave phenomenon. Especially, there is a lack of consistent experimental data for developing a three-dimensional constitutive model that predicts the essential feature of failure wave, and that could be easily implemented into a numerical code for large-scale computer simulation. Also, the relationships among different wave structures are still not clear.
Ceramic Armor Materials by Design 329
To the extent authorized under the laws of the United States of America, all copyright interests in this publication are the propertyof The American Ceramic Society. Any duplication, reproduction, or republication of this publication or any part thereof, withoutthe express written consent of The American Ceramic Society or fee paid to the Copyright Clearance Center, is prohibited.
The shock response of glasses beyond the apparent HEL often displays a distinctive two-wave structure in wave profile. The trailing longitudinal stress wave is referred as the inelastic shock wave. However, in the shock wave experiments reported so far, no obvious jump in the longitudinal stress history has been detected at the failure wave front.1, 2, 4, 9 One interpretation is that the apparent HEL may not be a true elastic limit, rather the manifestation of a transition in failure mechanisms. A possible transition is the one from a delayed kinetic-controlled failure process below the HEL to a prompt stress-controlled failure process above the HEL.12 Another possibility is that the HEL may represent the stress level above which bulk glass undergoes permanent densification.9 From existing experimental data, however, it appears that the signature feature that separates the failure wave from the usual inelastic shock wave is that only the lateral stress history is changed significantly while the longitudinal stress history remains almost constant at failed (due to the loss of shear strength) material particles. In other words, the propagation of a “failure wave” might not be the result of momentum balance. The underlying mechanism might be a process governed by a field equation other than the stress wave equation. Recent experiments conducted by Dandekar7 revealed that the longitudinal stress measured on the impact surface of a shocked glass plate is different from the stress measured at some distance from the impact surface during the propagation of the failure wave. From this observation, therefore, the formation and propagation of failure waves appear to depend not only on the local state, but also on the information in the domain of influence, which is similar to localization problems. In other words, a nonlocal approach should be considered to describe the failure wave. It has been demonstrated that the initiation and evolution of localized material failure can be related to the transition between governing field equations of different types.17, 18 Representing a hyperbolic-to-elliptic transition with a parabolic (diffusion) equation and using a local elastoplasticity model, Xin and Chen19 obtained an analytical solution for a dynamic softening bar. A diffusing failure front could be simulated via the jump forms of conservation laws, together with a local elastodamage model.5 Changes of governing equation type also arise in many thermal and fluid mechanics problems. For example, depending on the ratio of thermal diffusivity to relaxation time, heat may propagate at a finite speed as a thermal wave or at an infinite speed (in the absence of relaxation) as a thermal diffusion.20, 21 Two different elliptic equations may hold respectively inside and outside of a turbulence domain.22 Hence, multi-physics as reflected through the transition between governing equations is not unusual. From the available experimental results on the failure wave phenomenon, an attempt has been made to construct a micromechanics-based picture for the evolution of failure waves.11 It has been proposed that under plane shock wave
330 Ceramic Armor Materials by Design
loading, the material failure below the HEL occurs through simultaneousprocesses of heterogeneous microcracking, shear dilatancy and void collapsingunder high confining stresses, which result in an increase in the mean stress and a decrease in the deviatoric stress while all the longitudinal field variables remainunchanged. This particular form of failure initiates at the impact surface where thesurface defects and transient loading conditions are conducive for such a process,and propagates into the material bulk through progressive multiplication of microcracks, i.e., a percolation process with a certain threshold.
To develop an effective simulation procedure, a three-dimensional isotropiccontinuum damage model has been proposed based on the above micromechanics-based picture.23 The progressive percolation of micro-damage is described as a nonlinear diffusion process lagging behind the shock compression.Material dilatancy induced by shear microcracking is assumed and used toquantify the average intensity of damage. A unique feature of the proposed model is the postulation that the deviatoric elastic strain energy in the intact material is converted into the volumetric potential energy in the comminuted and dilated material during the time- and space-dependent failure evolution process. The twofield equations governing the elastic shock wave and the trailing damage diffusionare solved numerically via a staggered manner in a single computational domain.It appears that the simulations based on the proposed model and solutionalgorithm can predict the essential features of the stress histories associated withthe failure wave phenomenon as observed in plane shock wave experiments onsingle glass plates, with an assumed threshold. However, there is a lack ofunderstanding of the multi-physics and multi-scale effects on the initiation andevolution of dynamic structural failure. Especially, model parameters need becalibrated via consistent experimental data, and the change in the longitudinalstress profile, as observed in the experiments conducted by Dandekar,7 must be considered in the failure wave modeling.
Based on the recent experimental data of an aluminum target impacting on sandwiched glass plates, a numerical study is conducted here to simulate thelongitudinal stress histories measured at the copper-glass interface and the glass-glass interface, respectively.
CONSTITUTIVE MODELING AND DAMAGE DIFFUSIONFor the purpose of simplicity, a nonlinear elastic–perfectly plastic model is
used for metals, with an associated flow rule. The yield surface takes the form of
with 0sJ3f 2y2
pddss :
2
1J 2 , denoting the deviatoric stress tensor,
and being the yield strength. It can then be found that in the deviatoric space
ds
ys
Ceramic Armor Materials by Design 331
ddd
de
ss
ssPs d:
:G2
dddd (1)
with Pd denoting the deviatoric orthogonal projector, and ed being the deviatoricstrain tensor. Based on the shock physics, the pressure-dependent shear modulusis given by
200
m0
c
b41GG (2)
in which G is the original shear modulus, represents the initial mass density,
and b and can be determined via the relationship between the shock wave
velocity and particle velocity U , i.e., U . In the volumetric
space, the mean compressive stress is related to the current specific volume V
through the following equation:
0
c
sU
0
0
p p0s bUc
14exp4 0
0200
V
VVb
b
cm (3)
with V0 being the initial specific volume. The material parameters for copper havethe following values: 0=8924kg/m
3, b=1.51, c0=3910m/s, sy=60Mpa, and G0=49GPa, while for aluminum the values are 0=2703kg/m
3, b=1.34,c0=5350m/s, sy=40Mpa, and G0=25GPa.
To be complete, the essential ideas of the previous constitutive model forfailure waves in shocked glasses 23 are summarized here, with an emphasis on the modifications made. The diffusion equation governing damage variable Vd in the3-D space x with time t can be written as
dd Vtt
V,xD (4)
where D(x,t) denotes the second order damage diffusivity tensor. If themicroscopic details of percolation in different orientations are not pursued, it isreasonable to let with i being the second order identity tensorand
ixxD tDt ,,
332 Ceramic Armor Materials by Design
0
0
orif0
andif0,x
ddTHD
FHEL
F
ddTHD
VVYYYY
YYd
VVYY
tD (5)
where d is the diffusion coefficient, Y is the 2nd invariant of deviatoric stress, andsubscripts “F”, “HEL” and “THD” denote the values of the stress deviatorvariable in the completely failed but compressed material, at the HEL, and at thethreshold for initiating the failure process, respectively. Note that Y<YF during unloading.
It is assumed that the initial distribution of isolated microdamage sites near theimpact surface is sufficiently uniform and planar. As a result, the evolution offailure in a uniaxially shock-compressed material can be considered as a pseudo-3D process, in which the propagation of failure is described by longitudinaldiffusion supplemented with a time-dependent evolution function accounting forthe much faster lateral percolation of microdamage. Thus, we may approximateEqs. (4) and (5) by using
txQx
VtxD
xt
V dd ,, with 0,
, 0
d
dd
T
VV
d
txDtxQ (6)
where is the threshold below which is inactive, and T is the
characteristic time for the lateral evolution of microdamage at a givenlongitudinal location. The volumetric and deviatoric responses are modified to be
0dV txQ , d
m=45.36 e-137.0 e2 +208.3 e
3 and G=G0/[1+( x/ G)2], respectively, in unit of
GPa, with e=(V0+Vd)/V-1, x being the longitudinal stress, and G=5.0GPa and G0=30.43GPa for the glass material considered here.
SIMULATION AND DISCUSSIONThe experimental arrangement is shown in Fig.1. The tested material is a soda
lime glass with the density of 2450kg/m3 and the longitudinal sound speed of5.58km/s. The shock compression pulse was created by the impact of analuminum flyer plate which was launched by an explosive facility with a velocityof 1.17 0.05km/s. A series of numerical calculations based on the above materialmodels have been carried out to determine the material parameters that result inthe best simulation of the failure evolution characteristics as observed in theshocked sandwiched glass plates. Using Y=|s11-s22| as the stress deviator measure
Ceramic Armor Materials by Design 333
Fig. 1. Configuration of impact experiment.
Str
ess
(GP
a)
Time (µs)
simulated
measured
Fig. 2. Time histories of simulated and measured longitudinal stresses at copper-glass and glass-glass interfaces, respectively.
Time (µs)
lateral
longitudinal
S
tres
s (G
Pa)
Fig. 3. Time histories of simulated longitudinal and lateral stresses atcopper-glass and glass-glass interfaces, respectively.
334 Ceramic Armor Materials by Design
leads to the following values of the material parameters: YHEL=4.529GPa, YTHD=4.34GPa, and YF=0.75GPa for glass plate-1, and YHEL=4.529GPa, YTHD=3.76GPa, and YF=0.75GPa for glass plate-2. The parameters for damage diffusion in glass are of the following values: Vd0=5.0x10-7m3/kg, d=0.4m2/s, and Td=20ns. By assuming that each surface of the glass plates introduces additional microcracking sources as reflected through YTHD, the simulation results appear to match the available experimental data very well in the loading phase. However, the interactions among different kinds of waves in the unloading phase are still not clear from the experimental data available. An integrated experimental, analytical and computational study is needed to better understand the physics behind the impact failure responses of both single and composite materials.
ACKNOWLEDGMENT
This work was sponsored in part by National Science Foundation.
REFERENCES 1Brar, N.S., Bless, S.J., Rosenberg, Z., Impact-Induced Failure Waves in
Glass Bars and Plates. Applied Physics Letter 59, 3396-3398, 1991. 2Kanel, G.I., Rasorenov, S.V., Fortov, V.E., The Failure Waves and
Spallations in Homogeneous Brittle Materials. Shock Compression of Condensed
Matter–1991 (Edited by Schmidt, S.C., Dick, R.D., Forbes, J.W., Tasker, D.G.), Elsevier, 451-454, 1992.
3Bless, S.J., Brar, N.S., Impact Induced Fracture of Glass Bars. High-Pressure
Science and Technology (Edited by S.C. Schmidt, J.W. Shaner, G.A. Samana and M. Ross). AIP, New York, NY, 1813-1816, 1994.
4Bourne, N.K., Rosenberg, Z., Field, J.E., High-Speed Photography of Compressive Failure Waves in Glasses. Journal of Applied Physics 78, 3736-3739, 1995.
5Chen, Z., Xin, X., An Analytical and Numerical Study of Failure Waves. International Journal of Solids and Structures 36, 3977-3991, 1999.
6Clifton, R.J., Analysis of Failure Waves in Glasses. Applied Mechanics
Reviews 46, 540-546, 1993. 7Dandekar, D.P., Index of Refraction and Mechanical Behavior of Soda Lime
Glass under Shock and Release Wave Propagation. Journal of Applied Physics
84, 6614-6622, 1998. 8Dandekar, D.P., Beaulieu, P.A., Failure Wave under Shock Wave
Compression in Soda Lime Glass. Metallurgical and Material Applications of
Shock-Wave and High-Strain-Rate Phenomena (Edited by L.E. Murr, K.P. Staudhammer and M.A. Meyers). Elsevier Science B.V., 211-218, 1995.
Ceramic Armor Materials by Design 335
9Espinosa, H.D., Xu, Y., Brar, N.S., Micromechanics of Failure Waves in Glass: Experiments. Journal of the American Ceramic Society 80, 2061-2073, 1997a.
10Espinosa, H.D., Xu, Y., Brar, N.S., Micromechanics of Failure Waves in Glass: Modeling. Journal of the American Ceramic Society 80, 2074-2085, 1997b.
11Feng, R., Formation and Propagation of Failure in Shocked Glasses. Journal
of Applied Physics 87, 1693-1700, 2000. 12Grady, D.E., Shock-Wave Properties of Brittle Solids. Shock Compression
of Condensed Matter–1995 (Edited by Schmidt, S.C. and Tao, W.C), AIP, 9-20, 1996.
13Partom, Y., Modeling Failure Waves in Glass. International Journal of
Impact Engineering, 21, 791-799, 1998. 14Raiser, G., Clifton, R.J., Failure Waves in Uniaxial Compression of an
Aluminosilicate Glass. High-Pressure Science and Technology (Edited by Schmidt, S.C., Shaner, J.W., Samana, G.A., and Ross M.), AIP, 1039-1042, 1994.
15Raiser, G.F., Wise, J.L., Clifton, R.J., Grady, D.E., Cox, D.E., Plate Impact Response of Ceramics and Glasses. Journal of Applied Physics 75, 3862-3869, 1994.
16Rosenberg, Z., Bourne, N.K., Millett, J.C.F., Direct Measurements of Strain in Shock-Loaded Glass Specimens. Journal of Applied Physics 79, 3971-3974, 1996.
17Chen, Z., Continuous and Discontinuous Failure Modes. Journal of
Engineering Mechanics 122, 80-82, 1996.18Chen, Z., Sulsky, D., A Partitioned-Modeling Approach with Moving Jump
Conditions for Localization. International Journal of Solids and Structures 32,1893-1905, 1995.
19Xin, X., Chen, Z., An Analytical Solution with Local Elastoplastic Models for the Evolution of Dynamic Softening. Int. J. Solids and Structures 37, 5855-5872, 2000.
20Tzou, D.Y., On the Thermal Shock Wave Induced by a Moving Heat Source. ASME Journal of Heat Transfer 111, 232-238, 1989.
21Tzou, D.Y., Macro- to Microscale Heat Transfer: The Lagging Behavior,Taylor & Francis, Washington, DC, 1997.
22Chen, Z., Clark, T., Some Remarks on Domain-Transition Problems. Archives of Mechanics 47, 499-512, 1995.
23Chen, Z., Feng, R., Xin, X., Shen, L., A Computational Model for Impact Failure with Shear-Induced Dilatancy. Submitted for publication in International
Journal for Numerical Methods and Engineering, 2001.
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GRAIN LEVEL ANALYSIS OF CERAMIC MICROSTRUCTURES
SUBJECTED TO IMPACT LOADING
Pablo D. Zavattieri and Horacio D. Espinosa
Mechanical Engineering
Northwestern University
Evanston, IL 60208
ABSTRACT
A study on the accuracy of cohesive models for capturing dynamic
fragmentation of ceramic microstructures is presented. The investigation consists
of a combined experimental/numerical approach in which microcracking and
damage kinetics are examined by means of plate impact recovery experiments.
The numerical analysis is based on a 2-D micromechanical stochastic finite
element analysis. The model incorporates a cohesive law to capture microcrack
initiation and evolution as a natural outcome of the calculated material response.
Normal plate impact velocity histories are used not only to identify model
parameters, but also to determine under what conditions the model captures
failure mechanisms experimentally observed. The analyses show that in order to
capture damage kinetics a particular distribution of grain boundary strength and
detailed modeling of grain morphology are required.
INTRODUCTION
Critical elements in the development of a physically-based model of the
dynamic deformation and failure of ceramics requires experiments specically
designed to examine inelasticity. For instance, to study the initiation and
evolution of microcracks in ceramics, an experiment that can cause controlled
microcracking, under well defined stress conditions, was developed by Clifton
and co-workers [1, 2]. These investigators performed plate impact soft recovery
experiments by subjecting the central region of a square ceramic specimen to
known and controllable stress pulses. Microcracking resulted yet the specimens
were recovered intact for microscopic analysis.
A large portion of the microcracks was found to originate at triple points and
both inelasticity in compression and tension was interferometrically measured. In
the tension dominated region, several microcracks linked together to form a spall
Ceramic Armor Materials by Design 349
To the extent authorized under the laws of the United States of America, all copyright interests in this publication are the propertyof The American Ceramic Society. Any duplication, reproduction, or republication of this publication or any part thereof, withoutthe express written consent of The American Ceramic Society or fee paid to the Copyright Clearance Center, is prohibited.
plane perpendicular to the impact direction. In spite of these contributions to the
field of damage, lack of consensus on the mechanisms responsible for ceramic
failure under multi-axial dynamic loading still remains. Attempts have been made
to model the inelastic constitutive behavior of ceramics in the presence of cracks,
and to validate the models through simulation of plate and rod impact
experiments. Available models for the failure of ceramics are continuum damage
theories which are based on homogenizing the cracked solid and finding its
response by degrading the elasticity of the material, and discrete approachs [3, 4],
able to nucleate cracks, and follow their propagation and coalescence during the
deformation process, the influence of microscopic heterogeneities on the overall
material behavior, which depends on morphological characteristics such as size,
shape, lattice orientation and spatial distribution of grains, is not accounted for.
In order to provide powerful tools to understand the mechanisms that lead to
macroscopic failure and, at the same time, refine the theories of damage utilized
in continuum or continuum/discrete models, a 2-D micromechanical model is
presented to assess intergranular microcrack initiation and evolution. A
representative volume element (RVE) of an actual microstructure, subjected to
multi-axial dynamic loading, is considered for the different analyses. A large
deformation elastic-anisotropic visco-plasticity model for the grains,
incorporating grain anisotropy by randomly generating principal material
directions, is included. Cohesive interface elements are embedded along grain
boundaries to simulate intergranular fracture through microcrack initiation and
evolution. Their interaction and coalescence are a natural outcome of the
calculated material response.
This micromechanical model provides explicit account for arbitrary
microstructural morphologies and microscopic fracture patterns making it easier
to identify and design microstructural configurations that enhance fracture
toughness, and therefore lead to improvements in the manufacturing of ceramic
materials. A detailed study of the damage initiation and kinetics in soft-recovery
experiments is carried out.
The objective of this work is to provide tools and means to understand the
macroscopic inelastic response of ceramics when subjected to dynamic multi-
axial loading at the micron scale. This bridging between scales is achieved by a
micro-mechanical stochastic finite element model. Experiments are not only used
to examine and validate the micromechanical model but also to explain the
different failure mechanisms.
MICROMECHANICAL MODEL
The finite element analysis of the initial boundary value problem is performed
using a total Lagrangian continuum approach with a large deformation elastic and
thermal anisotropic visco-plastic model [5, 6]. The elastic and thermal anisotropic
350 Ceramic Armor Materials by Design
model is used to describe grains' single crystal anisotropic behavior. Each grain is
assumed to be elastic orthotropic and the orientation of the principal material
directions differs from grain to grain.
(a)
(b)
Figure 1: (a)Schematics of microcracking at grain boundaries using an irreversible
interface cohesive law.w (b) Soft-recovery normal impact configuration.
Ceramic Armor Materials by Design 351
A multi-body contact-interface algorithm is used to describe the kinematics at
the grain boundaries and to simulate crack initiation and propagation. Figure 1
describes the contact model, integrated with interface elements to simulate
microcracking at the grain boundaries and subsequent large sliding, opening and
closing of the interface. The tensile and shear tractions in the zero thickness
interface elements, embedded along grain boundaries, are calculated from the
interface cohesive law. The interface cohesive law describes the evolution of
these tractions in terms of both normal and tangential displacement jumps. Within
the framework of cohesive interface elements the two most noteworthy cohesive
failure models available in the literature are the potential-based law [7], and the
linear law [3]. More detail on the cohesive model used in this work can be found
in the following references [5, 6].
SOFT-RECOVERY IMPACT EXPERIMENTS
The “soft-recovery” plate impact experiment has been described in detail by
Raiser et al. [1], and Espinosa et al. [2]. The experiment uses an eight pointed
start-shaped flyer plate that impacts a square ceramic specimen, subjecting the
central octagonal region to a plane pulse. Figure 1(b), shows this soft-recovery
normal impact configuration. A tensile pulse is originated from a gap between the
specimen and the momentum trap upon reflection of the compressive pulse. The
velocity-time profiles recorded at the rear surface of the momentum trap plate
provide information on microcrack initiation and evolution.
Let x denote the distance from the front surface of the specimen measured in
the direction of impact, and let Ls denote the thickness of the specimen, Lf the
thickness of the flyer and LMT the thickness of the momentum trap. The particle
velocity induced in the rear surface of the momentum trap is measured as a
function of time by a normal displacement interferometer (NDI) and a normal
velocity interferometer (NVI).
In the case of brittle materials readily damaged in tension, the tensile region
becomes the likely site of substantial damage called spall region. When spallation
initiates, the release waves emitted from the newly created free surfaces
completely change the pattern of waves inside the specimen. The shape of the
pull-back signal and second compressive pulse reflects both microcracking, under
the tensile pulse itself, and attenuation while traveling through material already
damaged. The above one-dimensional analysis is valid in the central region of the
specimen, where the effects of diffracted waves from the corners and the edges of
the flyer are minimized [2]. The experimental findings suggested that the
modeling of crack nucleation and growth requires consideration not only of the
amplitude of the applied stress but also of its time dependence [2]. Several
successful tests have been conducted using this experimental design by Espinosa
352 Ceramic Armor Materials by Design
et al. [2] and Raiser et al. [1]. A summary of the shots used for comparisons with
the proposed numerical model can be found in Zavattieri and Espinosa, 2001 [8].
STOCHASTIC FEM SIMULATIONS
A representative volume element of an actual microstructure is considered for
the analysis. Although the exact grain geometry can be taken from a digital
micrograph, it is well established that the grain structure in polycrystalline
materials can be simulated by a Voronoi tessellation [6]. We followed the last
approach to generate enough statistical data.
Figure 2(a) shows a strip of the various plates used in the experimental
configuration, only the flyer, momentum trap and specimen are considered in the
analysis and due to the limited spread of tensile damage observed experimentally,
only a small portion of the ceramic in the spall region is simulated. The top and
bottom boundaries of the cell are modeled using viscous boundary conditions
which represent the exact elastic wave solution along characteristic lines. Details
on the boundary conditions and convergence can be found in [6].
ANALYSIS OF THE SOFT-RECOVERY IMPACT EXPERIMENTS
In order ot simulate these experiments, two important features were
incorporated in the simulation of the experiments, namely, (1) a Weibull
distribution of the interfacial strength and fracture toughness along the grain
facets. (1) Realistic microstructures considering grain morphology and size
distributions.
As discussed in [5, 6, 8], it is physically incorrect to select a uniform Tmax and
KIC for all grain facets. Not only that grain misorientation affect the interfacial
strength, but also it affect the presence of glassy phase, glass pockets, and other
impurities that modify grain boundary properties. In the following analyses, the
interfacial strength parameters will be described by a Weibull distribution.
Ceramic Armor Materials by Design 353
(a)
(b)
Figure 2. (a) Schematics of the computational cell used for the analyses. (b)
Experimental particle velocity versus time for one (Shot 88-04) of the
experiments performed by Espinosa et al. [2].
354 Ceramic Armor Materials by Design
Simulation of experiment 88-04
Figure 2(b) shows the experimental particle velocity vs time for shot 8804
performed by Espinosa et al. [2]. The impact velocity used in shot 88-04 was V0 =
48.4m/s. Figure 2(b) shows the experimental velocity history for this experiment.
The elastic solution is also shown in the same figure. The most significant
features of this experiment is the pullback signal (almost 30% of the maximum
stress) and the spreading of the second compressive pulse. Numerical simulations
using the microstructure shown in Figure 2(a) result in a pull-back signal with a
maximum stress equal to the first compressive pulse, which is well above the pull-
back signal measured experimentally. Once microcracks nucleate, they grow at
rates such that a major crack from side to side of the RVE develops. It is worth
noticing that even for the case in which there is only one nucleation site every
200µm, the crack has to propagate to the other side of the RVE, in more than 67
nanoseconds (tensile pulse duration), in order to have a pullback signal below
100% of the compressive pulse. This would require a crack speed of less than
50% of the Rayleigh wave speed, which for alumina is 3 mm/µs, or a delay in the
decohesion process produced by rate effects. From the SEM Micrographs [2], it is
observed that the microcracks need to follow grain boundaries, with large
variations in grain size. The net effect is that crack propagation speed on a
projected horizontal plane is reduced to a fraction of the Rayleigh wave speed.
We closely examine this feature in conjunction with the observation of possible
nucleation sites as a function of overstress from the threshold level.
Two microstructures are considered in this analysis. Both meshes have a
width of 300 µm such that if there is only one nucleation site, the crack will have a
total time equal to the pulse duration to coalesce into a main crack. The main idea
of this analysis is to compare vis-à-vis the crack propagation for two different
types of microstructures: Microstructure A, with a non-uniform distribution of
grain sizes and shapes (motivated from the micrographs), and microstructure B
with a uniform distribution of grains (all with the same size and similar shape).
Figure 3(a) shows in detail the pullback signal for simulations considering
microstructure A. Microstructures A and B are shown in Figure 4. In these
simulations three different Weibull distributions have been considered. The best
fit is obtained for a Weibull distribution with T = 5 GPa, = 2 MPa · m0
max
0
ICK1/2
and m =3. This distribution contains interface elements with Tmax = 0.5 GPa and
Tmax 10GPa. The same distributions have been considered for microstructure B,
see Figure 3(b), and the pullback signals are much more pronounced than those
obtained with microstructure A. An explanation can be inferred by examining the
evolution of crack patterns as shown in Figure 4. The evolution of the
microcracks for T = 5 GPa, = 2 MPa · m0
max
0
ICK1/2
and m =3 using mesh A is
shown in Figure 4(a); the grain morphology is shown in the first frame. In this
Ceramic Armor Materials by Design 355
(a)
(b)
Figure 3: (a) Comparison between three different Weibull distribution for shot 88-
04 using mesh A. (b) Comparison between the velocity history using meshes A
and B.
356 Ceramic Armor Materials by Design
(a)
(b)
Figure 4: (a) Evolution of crack pattern for the case with T = 5GPa and m =3
using mesh A (b) Evolution of crack pattern for the case with T = 5GPa and
0
max
0
max
m =3, using mesh B.
Ceramic Armor Materials by Design 357
case, it can be observed that the microcracks need to go around the large grains at
the center of the RVE. The time that it takes the crack to surround the large grains
is similar to the pulse duration and then the pullback signal is significantly lower
than for cases where the crack propagates, from one side to the other of the RVE,
at uniform speed. Figure 4(b) shows the crack evolution for the case with
microstructure B. The crack initiates almost in the middle of the RVE and
propagates at constant speed until it coalesces into a main crack just before the
tensile pulse vanishes. As a result, the pullback signal for this case is much higher
than that for the case where the crack is forced to follow a path around large
grains.
Higher impact velocities
In this subsection we examine an experiment with higher impact velocity. The
experiment (shot 92-11) has been reported by Raiser et. al [1], the initial velocity
was V0 = 92.3m/s. The main variation in this experiment is the average grain size
of the ceramic, Coors AD-999, of approximately 3 µm. For this analysis an RVE
of 200 x 200µm is considered and two type of microstructures, uniform and bi-
modal grain sizes, are analyzed. The main motivation for examining two different
microstructures is to study the effect of grain morphology and how this affect the
crack path and crack speed along the spall plane. Although the second
microstructure with a bi-modal distribution of grain sizes may not be totally
representative of the tested ceramic, it is used to evidence the effect of grain
morphology.
An analysis with three different Weibull distribution on the RVE with uniform
grain size has been carried out; weak interface case: T = 3GPa and m =3; the
case considered in previous experiments, i.e., T = 5GPa and m =3; and a strong
interface case: T = 10GPa and m =10. The intention of this analysis is not to
study parametrically the effect of m, or T
0
max
0
max
0
max
0
maxT
max. In all cases KIC= 2 Mpam1/2
. Figure 5
shows the crack pattern for each one of these cases; the grain morphology is
shown in the first frame. In the weak interface case, the ceramic fails from side to
side right after the tensile pulse is generated at the spall plane. Crack nucleation
occurs basically at a large percentage of triple points and coalescence of
microcracks occurs before the end of the tensile pulse. For the case with T =
5GPa and m =3 the crack start propagating from the center to the borders and
crack branching in the form of a “funnel” is observed. As it is expected, the
strongest case ( = 10GPa and m =10) shows less branching and microcrack
density. The energy to create new surfaces is higher so that branching is inhibited.
0
max
358 Ceramic Armor Materials by Design
Figure 5: Velocity history for Shot 92-11 using a microstructure with a uniform
distribution of grain size and three different Weibull distributions.
DISCUSSION
The micromechanical analyses, together with the experimental velocity
profiles and SEM observations, have demonstrated that there are two factors to be
taken into account to capture the right damage kinetics occurring during the
experiments. In view that not all grain facets have the same interface strength and
local fracture toughness, it is important to consider Weibull distributions of Tmax
and KIC. Similarly, since the ceramic microstructures interrogated in these
experiments do not contain grains with the same shape and size, microstructures
with non-uniform distributions of grain size and shape must be considered. On the
other hand, microstructures with non-uniform distribution of grain size and shape
strongly affect crack speed along the spall plane.
From a computational standpoint, simulations of ballistic penetration, vehicle
crash analysis, manufacturing processes, etc. cannot be conducted at the grain
level. Hence, this fundamental study of brittle failure provides insight into the
Ceramic Armor Materials by Design 359
utilization of cohesive laws at other size scales. Our simulations clearly show that
the scale at which simulations are performed plays an important role in the
selection of cohesive models. The calculations in this work make assumptions that
limited the degree of achievable accuracy. For instance, the model is two-
dimensional and crack interaction is stronger than in the 3-D case and therefore,
the computed rate of crack coalescence may be thought of as an upper bound.
Despite these limitations, the numerical results obtained with this model were not
only in good agreement with the experiments, but also were used to explain
several microscopic failure mechanisms that have never been quantified before
through other mathematical models.
REFERENCES
1. Raiser G., Wise J.L., Clifton R.J., Grady D.E., and Cox D.E., “Plate impact
response of ceramics and glasses”, J. Appl. Phys., 75(8):3862-69,1994.
2. Espinosa H.D., Raiser G., Clifton R.J., and Ortiz M., “Experimental
observations and numerical modeling of inelasticity in dynamically loaded
ceramics”, J. Hard. Mat., 3:285-313, 1993.
3. Camacho G.T. and Ortiz M. “Computational modeling of impact damage in
brittle materials”, Int. J. Sol. Str., 33: 2899-2938, 1996.
4. Espinosa H.D., Zavattieri P.D., and Dwivedi S., “A finite deformation
continuum/discrete model for the description of fragmentation and damage in
brittle materials”, J. of the Mechanics and Physics of Solids, 46(10): 1909-1942,
1998.
5. Zavattieri P.D., Raghuram P.V., and Espinosa H.D., “A computational model
of ceramic microstructures subjected to multi-axial dynamic loading”, J. of the
Mechanics and Physics of Solids, 49(1): 27-68, 2001.
6. H.D. Espinosa and P.D. Zavattieri, “A grain level model for the study of
dynamic failure of polycrystalline materials. Part I: Theory and numerical
implementation”, Submitted to Mechanics of Materials, 2001.
7. Xu X-P and Needleman A., “Numerical simulation of dynamic interfacial
crack growth allowing for crack growth away from the bond line”, Int. J. Fra.,
74:253-275, 1995.
8. P. D. Zavattieri and H. D. Espinosa, “Grain level analysis of crack initiation
and propagation in brittle materials”, In press Acta Materialia, 2001.
360 Ceramic Armor Materials by Design
ANALYSIS AND MODELLING OF CERAMIC ARMOUR PENETRATION
S.J. Cimpoeru and R.L. Woodward†
DSTO Aeronautical and Maritime Research Laboratory,
P.O. Box 4331, Melbourne, 3001, Australia.
ABSTRACT
This paper summarises fragmentation and energy absorption studies
conducted on a wide range of ceramic armour materials. Aspects of ceramic
armour depth of penetration tests are examined and how such data can be
interpreted and depend on test configuration. The Woodward one-dimensional
momentum balance model is also outlined and is used to represent some of the
characteristics of depth of penetration tests, including the importance of the
destruction of the penetrator nose, erosion of the remainder of the projectile and
the derivation of effective ceramic strength values.
FRAGMENTATION AND ENERGY ABSORPTION STUDIES
Quantitative fragmentation studies and ballistic performance measurements
were made on ceramics of a wide range of hardnesses and toughnesses, including
soda lime glass and zirconia, and on ceramics with similar measured mechanical
and physical properties but different ballistic performances. An inverse
correlation was found between between the volume of fragments produced and
fracture toughness [1,2]. It is of interest that any variation in fragmentation
behaviour due to small differences in toughness was masked by the shot to shot
inconsistency in the results and that significant variations in fragmentation
occurred despite similar residual depths of penetration. However, no correlation
was established between toughness and ballistic performance, which was to be
expected as measurements of the surface area of fractured ceramics and
calculations of fracture work [3,1] demonstrated that very little of the initial
projectile kinetic energy goes into fracturing the ceramic. A large proportion of
this energy simply ends up as the residual kinetic energy of the ejected ceramic
debris [3,1].
A marked difference was found in the fragmentation behaviour of blunt and
pointed projectiles, but this depended on whether the hardness of the ceramic was
†Deceased
Ceramic Armor Materials by Design 361
To the extent authorized under the laws of the United States of America, all copyright interests in this publication are the propertyof The American Ceramic Society. Any duplication, reproduction, or republication of this publication or any part thereof, withoutthe express written consent of The American Ceramic Society or fee paid to the Copyright Clearance Center, is prohibited.
sufficient to blunt the projectile upon impact [1]. For pointed projectiles
impacting soft ceramics, where the projectile remains undeformed, the ballistic
performance was found to increase with ceramic hardness. For hard ceramics the
residual penetration depths were generally similar due to the ability of these
ceramics to blunt pointed projectiles. Thus the ballistic performance of the hardest
ceramics is not simply related to hardness, but more related to the velocity
reduction and erosive mass loss incurred by the blunted projectile.
A SIMPLE ONE-DIMENSIONAL APPROACH TO MODELLING CERAMIC
COMPOSITE ARMOUR DEFEAT
Woodward [4,5] developed two one-dimensional momentum balance models
to highlight the essential physical processes of ceramic armour defeat. One model
was for the perforation of targets with thin backings which under the influence of
the ceramic fracture conoid deform by dishing, i.e. stretching and bending
deformation. A second model, a simplification of the thin backing model, was
developed for thick backings where the backing remains stationary whilst the
ceramic is eroded, allowing direct analysis of depth of penetration test results.
A analysis of the early work of Wilkins [6] concluded that the resistance to
penetrator motion was initially determined by the inertia of the ceramic and
backing that was bounded by the conoidal cracking of the ceramic and dishing of
the backing plate [4], i.e. a target acceleration stage. A second failure stage of
perforation was also identified where the penetrator and target material bounded
by the cone, moved forward at a common velocity until they are either slowed to
zero velocity or the backing plate ruptures via biaxial tensile failure [4]. An
additional failure case is also where the ceramic is eroded to zero thickness and
the backing plate is perforated according to a simple failure criterion [7].
Figure 1 shows the lumped mass model for the thin backing model [4]. In any
time step, t, a mass MP and a mass MC are eroded from the projectile and
ceramic, respectively, if the interface forces exceed the strength of the projectile
or ceramic, i.e. FI either exceeds FP or FC or both. Once ceramic erosion has
occurred, the area of the backing that is loaded is reduced to account for the
reduced mass distribution of ceramic in the fractured conoid, with a consequent
reduction in energy absorption.
Figure 1: Basic concepts of the lumped mass model [4].
362 Ceramic Armor Materials by Design
The model predictions were originally validated against the experimental results
of a number of other workers [4,5] and further validations have occurred in recent
years, e.g. [8].
INTERPRETATION OF DEPTH OF PENETRATION TEST DATA
Depth of penetration testing has become a popular low-cost means of
evaluating the ballistic performance of the wide range of ceramic options that are
now available. However, there are some disadvantages to this form of testing,
including the interpretation of the results in terms of the fundamentals of
penetration mechanics and a clear understanding of what material characteristics
are being measured. There are also a number of commonly used weight and space
merit ratings; understanding the meaning and therefore the limitations of the
ratings is as important as the values of the ratings themselves [9].
Depth of penetration test results are configuration dependent. Table I lists
residual depth of penetration results and derived ballistic and penetration
efficiency values for tungsten alloy and hardened steel projectiles, impacting
laterally confined alumina tiles that have different backing materials [10].
Table I:Depths of penetration and efficiency values for 99.5% alumina-faced targets [10].
WP—pointed tungsten alloy; WB—blunt tungsten alloy; SP—pointed steel; SB—blunt
steel.
ProjectileTarget
Backing
Reference
Depth
(mm)
Residual
Depth
(mm)
Ballistic
Efficiency [11]
Penetration
Efficiency [12]
WP AL 265 66 16.9 3.4
WB AL 75 57 1.5 1.1
SP AL 85 28 4.8 2.1
SB AL 62 20 3.6 2.0
WP RHA 34 20 3.4 1.4
WB RHA 23.5 14 2.3 1.3
SP RHA 30 2.5 6.7 4.6
SB RHA 10 3 1.7 1.4
WP HHS 23.5 13.7 2.4 1.3
WB HHS 23 13 2.5 1.3
SP HHS 19 1.5 4.3 3.4
SB HHS 11.5 3 2.1 1.6
Figure 2(a) shows the most stark of the comparisons of Table I in the form
suggested by Rosenberg and Yesherun [11] for the determination of ballistic
efficiency values, which equates to the slope of such plots. The standard tungsten
round penetrates the aluminium reference target without deforming, whereas the
truncated round of similar velocity and mass is a much less efficient penetrator
Ceramic Armor Materials by Design 363
because it deforms and erodes during its penetration into the reference target.
While a marked difference is seen between the reference penetration depths of
pointed and blunt penetrators, the penetration depth of both rounds is similar
when fired through ceramic tiles, because the immediate effect of impact onto the
hard ceramic is the destruction of the nose of the standard, pointed round. Thus in
both cases a blunt penetrator of similar geometry effectively impacts the ceramic,
in one case the blunt penetrator is pre-formed, and in the other case it is formed
by impact. The use of both pointed and blunt projectiles therefore allows the
magnitude of the nose fracture effect to be separated from velocity reduction and
erosive mass-loss effects.
Figure 2: Plots of AlhAl against chc: (a) for tungsten alloy projectiles, pointed,
, and blunt, ■, fired into 99.5% alumina-faced, aluminium-backed targets [10]
and (b) for hardened steel projectiles fired into AD85 alumina-faced, aluminium-
backed targets [11]. The two calculated intercepts in (b) are joined to obtain a
straight, dashed line to compare with the empirical result .
(b)
AlhAl
(kgm-2
)AlhAl
(kgm-2
)
chc (kgm-2
)
Experimental
(a)
chc (kgm-2
)
Some progress has been made in understanding behaviour by exercising models
of penetration mechanics in comparison with the experiments. Figure 2(b)
compares the data of Rosenberg and Yesherun [11] with calculations which
required three models. The penetration through the ceramic was calculated using
the Woodward momentum balance model [4,5]. A separate deep penetration
model for deformable, blunt penetrators was used to calculate the small residual
penetration into the backing material [13,14], which also established the intercept
on the abscissa. The ceramic strength was taken as 5.6 GPa, which is less than the
364 Ceramic Armor Materials by Design
Vickers Hardness, 8.8 GPa, but chosen to give a reasonable correspondence with
the results of Rosenberg and Yesherun [11].
The depth of penetration into the reference target was calculated using a
model for ductile hole formation by a non-deforming projectile [7]. An alternative
calculation for reference penetration was made using the deep penetration model,
which assumes a blunt projectile, and as shown it gives a much smaller reference
penetration. This illustrates the large effect on results of blunting the projectile tip,
particularly when an aluminium backing is used, and a comparison between
Figures 2(a) and 2(b) shows that the magnitude of the blunting effect observed in
experiments can be predicted with simple models.
Figure 2(b) illustrates that care should be exercised in using the slopes of such
curves to estimate ballistic efficiencies. The straight lines are usually constructed
from one point for the reference target and a few points from ceramic faced
targets that have little residual penetration into the backing armour with no
substantial data set in between. The relationship is not necessarily linear,
particularly at applique areal densities that are near to preventing penetration into
the reference target.
Table I also shows why ballistic efficiency values should be critically
examined and understood. For example, the ballistic efficiency value of 16.9
obtained for pointed tungsten alloy projectiles against an aluminium backing,
should be compared to the corresponding value of 1.5 for a blunt projectile. While
there is a significant difference between these merit ratings, the real effect of the
applique is almost identical in terms of residual penetration capability. Another
example is that given that the relative density of steel and alumina is 2.0, it
follows that a ballistic efficiency close to this value would imply similar
resistance per unit thickness for steel and alumina targets as also shown by Senf,
et al. [12]. It is seen that a number of ballistic efficiencies are close to this value
for the steel targets listed in Table I. Clearly the ceramic has not achieved the full
potential of its strength, despite its much greater hardness.
DERIVED VALUES FOR CERAMIC STRENGTH
Perhaps the greatest uncertainty with the application of one-dimensional
models to ceramic armour penetration problems is the difficulty in obtaining a
unique easily measured parameter, related to material strength, that indicates
resistance to penetration. While the indentation hardness of the ceramic is not
only easily measurable but also a physically meaningful strength parameter, other
measures of strength may be more suitable. Sternberg [15], for instance, indicated
that the penetration resistance is initially governed by the indentation hardness but
then drops to some lower value when cracking precedes penetration, the strength
parameter possibly being affected by ceramic toughness and confinement. It has
also been suggested that the appropriate strength value is a function of velocity
Ceramic Armor Materials by Design 365
[16], being low at low velocities where cracking can preceed penetration, but
higher as velocity increases closer to the rate of propagation of the damage front
[17].
The momentum balance model developed by Woodward [4,5], and
independently for finite thickness targets by den Reijer [18], bears a close
similarity to the modified hydrodynamic theory of penetration [19-21] which is
given by the equations
(1a)p p
Lu
(1b)L p u
p p c cu p p
1
2
2 1
2
2 (1c)
where p and c are the penetrator and target densities, respectively, L is the
penetrator length, p the depth of penetration, u the penetrator velocity, and p and
c are the strengths of penetrator and target, respectively. Formulating the
Woodward momentum balance model in the same way leaves equations 1(a) and
1(b) as they are and the third equation becomes
(2) p p c c
u p u p2
The similarity between equations (1c) and (2) is such that the methods give
similar results with different (but proportional) values for target strength.
Figure 3: Derived ceramic strength data for various impact velocities and tile
thicknesses: 20 mm, ◆; 25 mm, ■; 30 mm, ; 40 mm, ; and 50 mm, .
0
1
2
3
4
5
0 5 10 15 20
Momentum Balance (GPa)
Modif
ied H
ydro
dynam
ic (
GP
a)
1250 ms-1
1000 ms-1
1500 ms-1
366 Ceramic Armor Materials by Design
Despite the success in using such one-dimensional models in Figure 2(b), it is still
informative to analyse the velocity dependent data of Senf et al. [12] on the
penetration of 99.5% alumina-faced semi-infinite RHA targets. Figure 3 compares
ceramic strength values required by the models to match the data at impact
velocities between 1000 and 1500 ms-1
, with the deep penetration model [13,14]
being used to calculate the residual penetration into the backing armour. The
strength values for the two models are proportional to each other over a range of
tile thicknesses, but should be compared to the hardness of the alumina, 12.6 GPa
[12], and its HEL, 6.67 GPa [22]. Importantly, the proportionality is also found to
depend on impact velocity. Calculations such as this indicate that the models are
to some extent incomplete and as such should be applied with care.
CONCLUSION
A general inverse correlation has been found between the degree of ceramic
fragmentation and fracture toughness. It has been found also that a negligible
proportion of projectile kinetic energy is converted into surface area while a large
proportion of this energy ends up as the residual kinetic energy of the ejected
ceramic debris. Depth of Penetration (DOP) testing and the associated efficiency
ratings do not measure a fundamental armour property because they are test
configuration dependent and are usually strongly influenced by the fracture of the
projectile nose. The use of DOP data for real armour designs therefore requires a
close correspondence between the test set-up and the expected service
configuration. Current one-dimensional analytical models are able to represent
some of the characteristics of DOP tests such as the approximate calculation of
ballistic efficiencies. However, accurate prediction of performance over a range of
impact conditions requires a better understanding of ceramic strength effects and
so at present such models should be used with care.
REFERENCES
1. R.L. Woodward, W.A. Gooch, Jr, R.G. O’Donnell, W.J. Perciballi, B.J.
Baxter and S.D. Pattie, “A Study of Fragmentation in the Ballistic Impact of
Ceramics,” Int. J. Impact Engng, 15, 605-618 (1994).
2. R.G. O’Donnell, “An Investigation of the Fragmentation Behaviour of
Impacted Ceramics,” J. Mat. Sci. Lett., 10, 685-688 (1991).
3. R.L. Woodward, R.G. O’Donnell, B.J. Baxter, B. Nicol and S.D. Pattie,
“Energy Absorption in the Failure of Ceramic Composite Armours,” Materials
Forum, 13, 174-181 (1989).
4. R.L. Woodward, “A Simple One-Dimensional Approach to Modelling
Ceramic Composite Armour Defeat,” Int. J. Impact Engng, 9, 455-474 (1990).
Ceramic Armor Materials by Design 367
5. R.L. Woodward, “A Basis for Modelling Ceramic Composite Armour
Defeat,” Materials Research Laboratory, Australia, Research Report MRL-RR-3-
89 (1989).
6. M.L. Wilkins, “Computer Simulation of Penetration Phenomena”; pp. 225-
252 in Ballistic Materials and Penetration Mechanics. Edited by R.C. Laible.
Elsevier, Amsterdam, 1980.
7. R.L. Woodward, “The Penetration of Targets by Conical Projectiles,” Int. J.
Mech. Sci., 20, 349-359 (1978).
8. R. Zaera and V. Sánchez-Gálvez, “Analytical Modelling of Normal and
Oblique Ballistic Impact on Ceramic/Metal Lightweight Armours,” Int. J. Impact
Engng, 21, 133-148 (1998).
9. S.J. Cimpoeru, B.J. Baxter and R.L. Woodward, “Some Extensions of
Simplified Ballistic Test Procedures to Comparative Protection Analysis,” Proc.
16th Int. Symp. on Ballistics, ADPA, San Francisco, CA, USA. 23-27 Sept. 1996,
3, 17-26 (1996).
10. R.L. Woodward and B.J. Baxter, “Ballistic Evaluation of Ceramics:
Influence of Test Conditions,” Int. J. Impact Engng, 15, 119-124 (1994).
11. Z. Rosenberg and Y. Yesherun, “The Relation Between Ballistic
Efficiency and Compressive Strength of Ceramic Tiles,” Int. J. Impact Engng, 7,
357-362 (1988).
12. H. Senf, E. Strassburger, H. Rothenhäusler, W.A. Gooch and M.S.
Burkins “Ballistic Resistance of AD995 Al2O3 Ceramics against Short Projectiles
at Impact Velocities Between 1000 and 2000 m/s,” Proc. 15th Int. Symp. on
Ballistics, ADPA, Jerusalem, Israel. 21-24 May 1995, 1, 361-376 (1995).
13. R.L. Woodward, “Penetration of Semi-Infinite Metal Targets by
Deforming Projectiles,” Int. J. Mech. Sci., 24, 73-87 (1982).
14. R.L. Woodward, “Modelling Penetration by Slender High Kinetic Energy
Penetrators,” Materials Research Laboratory, Australia, Report MRL-R-811
(1981).
15. J. Sternberg, “Materials Properties Determining the Resistance of
Ceramics to High Velocity Penetration,” J. Appl. Phys., 65, 3417-3424 (1989).
16. Y. Partom and D. Littlefield, “Dependence of Ceramic Armor Resistance
on Projectile Velocity”, Proc. 14th Int. Symp. on Ballistics, ADPA, Québec,
Canada. 26-29 Sept. 1993, 2, 563-572 (1993).
17. Hornemann, H. Rothenhäusler, H. Senf, J.F. Kalthoff and S. Winkler,
“Experimental Investigation of Wave and Fracture Propagation in Glass Slabs
Loaded by Steel Cylinders at High Impact Velocities”; pp. 291-298 in Mechanical
Properties at High Rates of Strain. Edited by J. Harding. Institute of Physics
Conference Series No. 70, Institute of Physics, Bristol and London, 1984.
18. P.C. den Reijer, “Impact on Ceramic Faced Armour,” Doctoral Thesis,
Delft University of Technology, The Netherlands, (1994).
368 Ceramic Armor Materials by Design
19. V.P. Alekseevkii, “Penetration of a Rod into a Target at High Velocity,”
Fiz. Goren. Vzryva, 2, 99-106 (1966).
20. A Tate, “A Theory for the Deceleration of Long Rods After Impact,” J.
Mech. Phys. Sol., 15, 387-399 (1967).
21. A. Tate, “Further Results in the Theory of Long Rod Penetration,” J.
Mech. Phys. Sol., 17, 141-150 (1969).
22. D.P. Dandekar and P. Bartkowski, “Strength of AD995 Alumina under
Impact Loading,” 1994 Army Science Conf., Orlando, Florida, June 1994.
Ceramic Armor Materials by Design 369
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OVERVIEW OF THE RAJENDRAN-GROVE CERAMIC FAILURE MODEL
D. J. Grove A. M. Rajendran
U. S. Army Research Laboratory U. S. Army Research Laboratory
APG, MD 21005-5067 ARO, RTP, NC 27709-2211
ABSTRACT
This paper provides an up-to-date detailed description of the Rajendran-Grove
(RG) ceramic failure model. Damage initiates and evolves when the stress state
satisfies either a generalized Griffith criterion or a spall criterion. The ceramic
material's stiffness decreases as the microcrack damage increases, and microcrack
coalescence is assumed to occur at a critical crack density. The RG model has
been permanently implemented in the latest version of the EPIC code. EPIC
simulations were performed to evaluate the predictive capabilities of the model
for a benchmark suite of dynamic impact experiments. Results from these
calculations are presented and discussed in this paper.
INTRODUCTION
Historically, constitutive damage models have not demonstrated the predictive
capabilities necessary to justify their widespread use in the design of armor/anti-
armor systems. Empirical models are relatively easy to use (i.e., small number of
constants, etc.) and computationally efficient, but they tend to be applicable only
for a limited set of loading conditions. Three-dimensional fracture mechanics
based microphysical models have more complex formulations that tend to require
larger numbers of model constants, some of which may be difficult to determine
from experimental measurements. In addition, since their solution algorithms
may include iterative procedures to produce accurate results, these models tend to
require significantly more computing resources to solve a problem. However, a
model formulation which addresses the microphysics of the damage evolution
process offers the greatest potential for achieving an accurate predictive
capability. This paper describes the Rajendran-Grove (RG) microphysical
ceramic failure model in detail, and then demonstrates its ability to reproduce and
predict experimental measurements from a variety of impact experiments on
99.5% pure aluminum oxide (AD995) ceramic targets.
Ceramic Armor Materials by Design 371
To the extent authorized under the laws of the United States of America, all copyright interests in this publication are the propertyof The American Ceramic Society. Any duplication, reproduction, or republication of this publication or any part thereof, withoutthe express written consent of The American Ceramic Society or fee paid to the Copyright Clearance Center, is prohibited.
MODEL DESCRIPTION
The Rajendran-Grove (RG) ceramic failure model (Rajendran [1], Rajendran
and Grove [2]) assumes the following: 1) preexisting randomly distributed flaws,
2) plastic flow and pore collapse when shocked above the Hugoniot elastic limit
(HEL), 3) no plastic flow in tension, 4) degradation of elastic moduli under both
compression and tension due to microcracking, and 5) pulverization at a critical
crack density. A strain-rate-dependent strength relationship is employed to
describe the response of the ceramic material due to inelastic (plastic) deformation
under high-compression pressure. The deviatoric stresses are calculated using a
conventional radial return approach that is often used in viscoplasticity models.
The total strain is decomposed into elastic and plastic strains. The elastic strains,
which consist of the elastic strains in the intact matrix material and the strains due
to crack opening/sliding, are obtained by subtracting the plastic strains from the
total strains. After damage initiates, the unloading and reloading paths follow the
degraded elastic modulus (secant modulus), thus allowing full recovery of the
strains due to microcracking. The elastic stress-strain relationship for the
damaged aggregate material is given by
pklklijklij M , (1)
where is the total stress rate, is the total strain rate, is the plastic
strain rate due to viscoplastic flow and pore collapse, the strain rate difference
is the elastic strain rate, and the components of the stiffness tensor M
are determined by analytically inverting the compliance tensor.
ij
)pkl
klpkl
( kl
For non-interacting, penny-shaped microcracks of various sizes and random
orientations, Margolin [3] provided the following expression for the elements of
an isotropic compliance tensor C:
C , (2)klijjkiljlikijkl ccc 321
whereG
Bo4
11c ,
GDo
4
12c , and
G12Ac o3 . (3)
In the above equations, is the Kronecker delta, G and are the shear modulus
and Poisson's ratio, respectively, of the undamaged material, and Ao, Bo, and Do
are parameters whose values depend on the principal stress state and the extent of
damage. Before damage initiates, Ao = Bo = Do = 0. For stress states in which all
three principal stresses are compressive, Margolin proposed that
372 Ceramic Armor Materials by Design
, , and , (4)*o 1A *
o 14B oo AD
whereE45
16* and . (5)3*
o aN
In the above equations,* is a microcrack density parameter, is the microcrack
density (a dimensionless measure of damage), E is the Young's modulus of the
undamaged material, (a model constant) is the number of microcracks per
unit volume, and a is the maximum crack length in the exponential distribution of
microcracks. Upon substituting Equations (4) into Equations (3) and inverting the
compliance tensor C, we obtain the following expressions for the shear and bulk
moduli (
*oN
cG and cK , respectively) of the damaged ceramic material under
compression:
GB32
G2
ocG (6)
and213
1G2Kc . (7)
The above equations result in degradation of the shear modulus, but not of the
bulk modulus, because under compression only crack movement of the closed
microcracks (modes II and III) is permitted. During tensile loading, the RG model
assumes that both the bulk and shear moduli rapidly degrade as a result of
microcrack opening (mode I). Based on the damaged stiffness solution proposed
by Budiansky and O'Connell [4] for randomly oriented, non-interacting
microcracks under tensile loading, we employ the following expressions for the
shear and bulk moduli ( tG and tK , respectively) of the damaged ceramic material
under tension:
t
ttt
2
51
45
321GG (8)
andt
ttt
213
1G2K , (9)
where9
161t . (10)
Ceramic Armor Materials by Design 373
The above equations (valid only for tension), limit the microcrack density to
9/16; at this critical value, the ceramic material experiences a complete loss of
stiffness ( tK = tG = t = 0) and is subsequently unable to carry any tensile load.
Microcrack growth/extension (increasing a) causes the microcrack density
to increase (see Equations (5)), which results in the relaxation of stresses in the
cracked ceramic material. The crack growth (damage evolution) law, derived
from a fracture mechanics based relationship for a single crack propagating under
dynamic loading conditions, is described by
2n
I
crR1
G
G1Cna , (11)
where C is the Rayleigh wave speed, G is the critical strain energy release rate
for microcracking, are the applied strain energy release rates, and and n
are model constants that are used to limit the microcrack growth/extension rates.
The "+" superscript corresponds to microcrack opening under tension (mode I),
while the "-" superscript relates to microcrack extension under compression
(modes II/III). The model constants , , and are always assumed to be
equal to 1.0, while n is generally assumed to be equal to 0.1. In the model,
microcracks nucleate and grow/extend (i.e., > 0) when the stress state satisfies a
generalized Griffith criterion [5] developed by Margolin [6] and Dienes [7]. For
this criterion, , G , and G are calculated as:
R cr
1n n
IG
1
I
1n 2
2
a
2n
crG I
E
1K 22IC
crG , (12)
kji,2
2
E
a14G
2jk
2ik2
kk
2
I , (13)
and kji,2E
a18 2
kk2jk
2ik
2
IG , (14)
where (a model constant) is the fracture toughness of the undamaged
material,
ICK
E and are the degraded Young's modulus and Poisson's ratio,
374 Ceramic Armor Materials by Design
respectively, of the damaged (cracked) material, and (a model constant) is the
dynamic friction coefficient. Note that G is computed from the undamaged
(virgin) material properties. Microcrack opening occurs when G exceeds ,
or microcrack extension occurs when exceeds .
cr
I crG
IG crG
lnC1
Y2
P3
m
f2f 2
Under high tri-axial tensile stress loading conditions, the following damage
processes often occur in brittle materials: 1) debonding of the hard carbide and
oxide particles from the matrix, and 2) non-spherical pore (planar crack) growth at
triple point grain boundaries. To capture the effects of damage due to these
processes, the RG model employs a critical stress based spall criterion (in addition
to the Griffith criterion). This spall criterion assumes initiation and growth of
damage when all three principal stresses are tensile and the maximum principal
stress exceeds a critical spall threshold stress, s (a model constant). The damage
rate in this case is assumed to be simply proportional to the Rayleigh wave speed
(i.e., ). The spall damage criterion is only applied to tensile stress states
that fail to satisfy the generalized Griffith criterion for crack opening/extension.
Consequently, the microcrack density is accumulated in a continuous manner
due to either microcracking or spall damage.
RCa
The RG model considers the material to be in a comminuted (pulverized) state
when the microcrack density ( ) exceeds p (a model constant) during compressive
loading. Generally p is set to 0.75, based on the assumption that pulverization
occurs when the microcracks coalesce [8].
Prior to pulverization, the compressive strength Y of the matrix (void-free)
ceramic material is described by the following strain rate dependent relationship:
peffAY , (15)
where A is the quasi-static maximum strength, C is the strain rate sensitivity
parameter, and is the normalized (dimensionless) equivalent plastic strain
rate; A and C are model constants. The model assumes that pore collapse may
occur during compressive loading above the HEL due to local microscopic plastic
flow in the matrix material surrounding the pores. The pore collapse strain
components are derived from Gurson's pressure dependent yield surface [9]:
peff
0cosh1YJ3 2m2 , (16)
where J2 is the second invariant of the deviatoric stress in the porous (void-
Ceramic Armor Materials by Design 375
containing) aggregate material, Ym is the effective stress in the matrix (void-free)
material, P is the compressive pressure in the porous aggregate material, and f is
the porosity (void volume fraction). Note that in the absence of porosity (i.e.,
when f = 0), Equation (16) reduces to the von Mises yield condition. If pore
collapse occurs, the effective shear and bulk moduli of the damaged aggregate
material are defined as follows using a modified form of Mackenzie's relationship
[10,11]:
oo
eff
fG8K9
G12K61f1
fG8K9
G12K61f1
GG (17)
and
fG4
K31f1
fG4
K31f1
KK
o
o
eff , (18)
where . (19)pvo ef11f
In the above expressions, and are the effective shear and bulk moduli
(respectively) of the aggregate material,
effG effK
G is the degraded shear modulus of the
cracked matrix material (either cG or tG ), K is the degraded bulk modulus of
the cracked matrix material (either cK or tK ), fo (a model constant) is the initial
porosity, (= ) is the plastic volumetric strain due to pore
collapse, f is the porosity, and G and K are the shear and bulk moduli
(respectively) of the virgin material.
pv
p33
p22
p11
The RG model employs the following modified Mie-Gruneisen relationship to
compute the pressure P in the aggregate material prior to pulverization:
0,1E5.01bK/K
0,1E5.01bbbK/KP
s1eff
s3
32
21eff
, (20)
where is the elastic volumetric compressive strain ( = / o - 1, where and o
are the current and initial densities, respectively), is the Gruneisen coefficient,
376 Ceramic Armor Materials by Design
b1, b2, and b3 are empirical constants used for a cubic fit to the Hugoniot curve for
the virgin material, Es is the internal energy per initial volume, K is the bulk
modulus of the virgin material, and is the effective bulk modulus of the
aggregate material (see Equation (18)).
effK
Once pulverization has occurred (i.e., p during compressive loading), the
RG model assumes that pore collapse and crack growth/extension no longer occur
(i.e., ). Also, the model assumes that the pulverized material is unable
to carry a tensile load, so that
0af
effK
ij = P = 0 in tension. In compression, however, the
stresses and pressure are computed using the effective shear and bulk moduli
( and ) corresponding to the values of and f at the time of
pulverization. The pressure P and strength Y of the comminuted material are
described by,
effG
(21)
0,0
0,KP
ev
ev
eveff
and
0P,0
0P,Y,Pmin maxpp
Y , (22)
where is the elastic volumetric strain, (a model constant) is the dynamic
friction coefficient for granular motion, and Y (a model constant) is the upper
limit on the compressive strength of the pulverized ceramic material. Since
experimental data for the fractured strength is generally either unavailable or
difficult to interpret, we usually set to "1" and calibrate Y to match the
measured penetration depths from projectile penetration experiments.
ev p
maxp
pmaxp
Generally, seven of the RG ceramic model constants require some calibration
with experimental data: strain rate sensitivity parameter (C), initial crack size
( ), microcrack number density ( ), dynamic friction coefficient ( ),
coefficient to limit the mode II/III crack extension rate ( ), critical spall stress
(
oa *oN
1n
s), and the maximum compressive strength of the pulverized material (Y ).
We employed the following set of model constants to describe the dynamic
response of AD995 ceramic in this study:
maxp
o = 3.89 g/cm3, G = 156 GPa, K = 231
GPa, b1 = 231 GPa, b2 = -160 GPa, b3 = 2774 GPa, = 2.3, A = 2.3 GPa, C = 0.2,
KIC = 3 MPa m , fo = 0.0, a = 1.5 x 10o-6
m, = 2.0 x 10*oN
maxp
-11 m
-3, = 0.60,
= 0.1, 1n s = 0.5 GPa, p = 0.75, = 1.0, and Y = 4.5 GPa. p
Ceramic Armor Materials by Design 377
MODEL RESULTS FOR AD995 CERAMIC
We verified the generality of the RG model constants for 99.5% pure
aluminum oxide (AD995) through computer simulations of the following four
impact configurations: 1) plate impact, 2) rod-on-rod impact, 3) graded-density
plate-on-rod impact, and 4) projectile penetration. The details of these
experimental configurations are provided in a companion chapter, "Historical
Perspective on Ceramic Materials Damage Models," by A. M. Rajendran.
The simulations were performed using the EPIC finite element code, modified
to include the RG ceramic failure model. EPIC is a well-established three-
dimensional production code that was initially developed in the early 1970's to
describe the response of solid materials to dynamic impact loading. Johnson,
Stryk, Holmquist, and Beissel [12] have described the details of this explicit
Lagrangian finite element code. To maintain the stability of the explicit finite
element solution, an iterative scheme based on a second-order diagonally implicit
Runge-Kutta method was employed in the RG model solution algorithm.
Plate Impact
Using EPIC's one-dimensional (1D) strain option, we obtained an initial
calibration of the model constants through simulations of two AD995 plate impact
tests: 1) a low velocity test (flyer thickness: 4 mm, target thickness: 8 mm, impact
velocity: 83 m/s) performed by Dandekar and Bartkowski [13], and 2) a high
velocity test (flyer thickness: 5 mm, target thickness: 10 mm, impact velocity:
1943 m/s) reported by Grady and Moody [14]. Figure 1 indicates that the
computed spall signals (profiles beyond point S) agree with the measured profiles.
Time ( s)
0.5 1.0 1.5 2.0 2.5
Axia
l S
tres
s (G
Pa)
0.00
0.05
0.10
0.15
0.20
0.25
0.30
Experiment
Model
S
(a) Impact velocity = 83 m/s
Time ( s)
0.5 1.0 1.5 2.0 2.5 3.0
Vel
oci
ty (
km
/s)
0.0
0.2
0.4
0.6
0.8
1.0
1.2
1.4
Experiment
Model
S
(b) Impact velocity = 1943 m/s
Figure 1. Comparison of computational results with plate impact data.
378 Ceramic Armor Materials by Design
Rod-on-Rod Impact
Using the two-dimensional (2D) axisymmetric geometry option in EPIC, the
rod-on-rod impact configuration was simulated for two different impact velocities
(0.175 km/s and 0.278 km/s). The AD995 striker rod was 5 cm long and 1.25 cm
in diameter (L/D = 4), while the AD995 target rod was 10 cm long and 1.25 cm in
diameter (L/D =8). Figure 2 compares the model-predicted stress histories with
those measured experimentally by Simha [15]. While the model does not exactly
match the data, the peak stress levels are reasonably close to the measurements.
Time ( s)
6 7 8 9 10
Str
ess
(GP
a)
0
1
2
3
4
5
Experiment
Model
(a) Impact velocity = 175 m/s
Time ( s)
6 7 8 9 10
Str
ess
(GP
a)
0
1
2
3
4
5
Experiment
Model
(b) Impact velocity = 278 m/s
Figure 2. Comparison of computational results with rod-on-rod impact data.
Graded-Density Plate-on-Rod Impact
Two-dimensional axisymmetric simulations of the graded-density plate-on-rod
impact configuration were performed for both unsleeved (bare) and sleeved
AD995 ceramic target rods. In both cases, the impact velocities were around
0.300 km/s. The flyer plate was modeled as a layered circular disk (diameter = 5
cm, thickness = 2.2 cm); a continuous finite element grid was employed in the
flyer plate to simulate a "perfect" bond between adjacent layers of material. The
target was modeled as a solid rod (diameter = 1.9 cm, length = 7.4 cm), while the
steel sleeve was modeled as a hollow rod (inner diameter = 1.9 cm, outer diameter
= 3.8 cm, length = 7.4 cm). Frictionless sliding was permitted between the inner
surface of the sleeve and the outer surface of the ceramic rod. Figure 3 compares
the model-predicted velocity histories with those measured experimentally by
Chhabildas, Furnish, Reinhart, and Grady [16]. As the figure indicates, the model
does an excellent job of predicting the peak velocity levels, as well as the
constant-velocity behavior (due to spallation of the ceramic rod near the free end).
Ceramic Armor Materials by Design 379
(a) Bare target rod,
impact velocity = 300 m/s
(b) Sleeved target rod,
impact velocity = 321 m/s
Figure 3. Comparison of computational results with graded-density plate-on-rod
impact data.
Time ( s)5 10 15 2
Vel
oci
ty (
km
/s)
0.00
0.05
0.10
0.15
0.20
0.25
0.30
Experiment
Model
Time ( s)5 10 15 2
Vel
oci
ty (
km
/s)
0.00
0.05
0.10
0.15
0.20
0.25
0.30
Experiment
Model
00
Projectile Penetration
For the two-dimensional axisymmetric simulations of the projectile
penetration configuration, the target was assumed to be an AD995 ceramic disk
(diameter = 15.24 cm) backed by a thick steel cylinder (diameter = 20.32 cm,
thickness = 12.7 cm); the ceramic disk was also radially confined by a steel ring
(inner diameter = 15.24 cm, outer diameter = 20.32 cm, thickness = 5.08 cm) that
was fixed to the surface of the steel cylinder. The projectile was modeled as a
tungsten rod (diameter = 0.787 cm, length = 7.87 cm) with an impact velocity of
1.5 km/s. Simulations of this configuration were performed for seven different
thicknesses (between 1.02 and 5.08 cm) of AD995 ceramic disks. The penetration
process was modeled through EPIC's erosion algorithm (using an erosion strain of
150%). Figure 4 compares the measured [17] and computed residual depths of
penetration (DOP) into the backup steel block versus the areal densities (mass per
unit area) of the ceramic disks. The straight line in this figure is a linear least-
squares fit to the experimental data. As Figure 4 indicates, the model-predicted
depths of penetration are consistent with the experimental measurements.
SUMMARY
The governing equations for the RG ceramic failure model were described in
detail, and a set of model constants for AD995 ceramic was proposed. This set of
constants was then employed in a series of finite element simulations for the
following benchmark suite of experimental impact configurations: plate impact,
380 Ceramic Armor Materials by Design
Areal Density (g/cm2)
0 5 10 15 20 25
Res
idual
DO
P (
cm)
0
1
2
3
4
5
6
7
8
Experiments
Model
Figure 4. Comparison of computational results with projectile penetration data.
rod-on-rod impact, graded-density plate-on-rod impact, and projectile penetration.
The simulation results demonstrated the model's ability to reproduce the
experimentally measured stress and velocity histories, as well as the DOP data.
While these results are very encouraging, it is important to continue evaluating the
model's predictive capability through simulations of more complex ceramic armor
impact configurations.
ACKNOWLEDGEMENTS
The authors greatly appreciate the funding support of Dr. Doug Templeton and
Krishan Bishnoi of TARDEC, Warren, MI. This work was supported in part by a
grant of HPC time from the DoD HPC Center at Aberdeen Proving Ground, MD.
REFERENCES1A.M. Rajendran, "Modeling the Impact Behavior of AD85 Ceramic Under
Multiaxial Loading," Int. J. Impact Engng., 15 (6) 749-768 (1994). 2A.M. Rajendran and D.J. Grove, "Modeling the Shock Response of Silicon
Carbide, Boron Carbide, and Titanium Diboride," Int. J. Impact Engng., 18 (6)
611-631 (1996). 3L.G. Margolin, "Elastic Moduli of a Cracked Body," Int. J. of Fracture, 22,
65-79 (1983). 4B. Budiansky and R.J. O'Connell, "Elastic Moduli of a Cracked Solid," Int. J.
of Solids and Structures, 12, 81-97 (1976).
Ceramic Armor Materials by Design 381
5A.A. Griffith, "The Phenomena of Rupture and Flow in Solids," Phil. Trans.
of Royal Soc. London, 221, 163-198 (1920). 6L.G. Margolin, "A Generalized Griffith Criterion for Crack Propagation,"
Eng. Fracture Mechanics, 19 (3), 539-543 (1984). 7J.K. Dienes, "Comments on 'A Generalized Griffith Criterion for Crack
Propagation', by L.G. Margolin," Eng. Fracture Mechanics, 23 (3), 615-617
(1986).8A.M. Rajendran, High Strain Rate Behavior of Metals, Ceramics, and
Concrete, Air Force Report WL-TR-92-4006, Wright-Patterson Air Force Base,
OH 45433-6533, April 1992. 9A.L. Gurson, "Continuum Theory of Ductile Rupture by Void Nucleation and
Growth; Part I: Yield Criteria and Flow Rules for Porous Ductile Materials," J.
Engr. Mat. Tech., 99, 2-15 (1977). 10
J.K. Mackenzie, "The Elastic Constants of a Solid Containing Spherical
Holes," Proc. Phys. Soc., 2, 63, (1950). 11
J.N. Johnson, "Dynamic Fracture and Spallation in Ductile Solids," J. Appl.
Phys., 52 (4), 2812 (1981). 12
G.R. Johnson, R.A. Stryk, T.J. Holmquist, and S.R. Beissel, Numerical
Algorithms in a Lagrangian Hydrocode, Report No. WL-TR-1997-7039, Wright
Laboratory, Eglin AFB, FL (1997). 13
D.P. Dandekar and P. Bartkowski, "Shock Response of AD995 Alumina";
pp. 733-736 in High-Pressure Science and Technology - 1993, Part 1. Edited by
S.C. Schmidt, J.W. Shaner, G.A. Samara, and M. Ross. AIP Press, New York,
1994.14
D.E. Grady and R.L. Moody, Shock Compression Profiles in Ceramics,
Sandia Report No. SAND96-0551, Sandia National Laboratory, Albuquerque,
NM. (1996). 15
C.H.M. Simha, High Rate Loading of a High Purity Ceramic - One
Dimensional Stress Experiments and Constitutive Modeling, Ph.D Thesis,
University of Texas, Austin, Texas. (1998). 16
L.C. Chhabildas, M. D. Furnish, W.D. Reinhart, and D.E Grady, "Impact of
AD995 Alumina Rods"; pp. 505-508 in Shock Compression of Condensed Matter
- 1997. Edited by S.C. Schimdt, D.P. Dandekar, and J.W. Forbes. AIP Press,
1998.17
P. Woolsey, "Residual Penetration Ballistic Testing of Armor Ceramics,"
Unpublished Work, U. S. Army Materials Technology Laboratory, Watertown,
MA (1991).
382 Ceramic Armor Materials by Design
Damage Evolution and Micromechanisms
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FAILURE PHENOMENOLOGY OF CONFINED CERAMIC TARGETS AND
IMPACTING RODS
Donald A. Shockey and A.H. Marchand
SRI International
333 Ravenswood Avenue
Menlo Park, CA 94025
S.R. Skaggs, G.E. Cort, M.W. Burkett, and R. Parker
Los Alamos National Laboratory
Los Alamos, NM 87545
ABSTRACT
The mechanism by which a long rod penetrates a steel-encased ceramic block
was sought by performing impact experiments at a range of velocities, and
examining the fracture and deformation in the recovered targets and impactors.
The key processes are the crushing of a small volume of ceramic adjacent to the
leading surface of the advancing penetrator, and the subsequent flow of the fine
fragments lateral to and then opposite the direction of attack. The results suggest
that nonconventional material properties such as the dynamic compressive failure
energy and the friction, flow and abrasive properties of the finely fragmented
material govern the penetration resistance of confined ceramics. This
understanding of penetration mechanisms can be used to guide development of
specialized tests and failure models to measure pertinent material properties and
to predict penetration behavior, respectively.
INTRODUCTION
When ceramic plates are used as overlays or incorporated as a layer within
conventional monolithic steel armor, the ballistic protection is significantly
enhanced [1]. This finding has encouraged the use of ceramics as a component of
advanced armor structures and has motivated researchers to identify materials and
structural configurations that maximize ballistic performance.
The design of current ceramic armors is based predominantly on empirical
ballistic performance data. Test firings using the threat of interest (long rod,
shaped charge jet, or small caliber ammunition) are conducted against armor
specimens in which ceramic component parameters, such as type of ceramic,
Ceramic Armor Materials by Design 385
Reprinted from International Journal of Impact Engineering, Vol 9, No 3, Shockey et al, “Failure Phenomenology of Confined CeramicTargets and Impacting Rods”, pp 263-275, copyright1990, with permission from Elsevier Science.
thickness, and spacing of plates, are varied independently and systematically. The
combination of parameters that produce maximum ballistic protection is
determined from test results and used to design the armor package. This
procedure is lengthy and expensive, and because the number of potentially
influential material and geometry variables is large, a comprehensive test matrix is
not practical and so it is doubtful that armor packages affording optimal protection
are attained.
A more efficient procedure is to combine experiments with computational
simulations of experiments. Computations of the ballistic behavior for various
impact conditions and target geometries indicate which target parameters are
important and suggest combinations of target parameters that will give favorable
performance. A limited matrix of ballistic experiments based on these guidelines
are performed. The results are used to modify the models and the code; then a
second generation of computational simulations is conducted and used to design a
second set of test firings. This procedure is repeated until an armor package with
acceptable performance is obtained. The reduction in the number of required test
firings reduces substantially the expense and time required to attain a suitable
armor, and the understanding gained promotes optimum armor design.
Implementation of the iterative computational/experimental procedure,
however, requires reliable models for the microfailure behavior of ceramics and
penetrators under penetration conditions. These models should be based on an
understanding of the failure phenomenology during penetration. In particular, the
material properties governing penetration resistance must be known.
Unfortunately, conventional material properties such as fracture toughness,
strength, and hardness correlate poorly with penetration behavior [2, 3],
suggesting that under the complex, high-rate, multiaxial load produced by the
penetrator one or more nonconventional material properties control penetration.
The goal of the work reported here was to establish the failure phenomenology
of confined ceramic targets and impacting long rods during penetration, and to
deduce the ceramic properties governing penetration resistance. Our approach was
to perform impact experiments on confined ceramic specimens at several
velocities to produce damage ranging from incipient to severe. Very early stages
of damage were studied by performing experiments with low velocity spherical
particles. From fractographic and metallographic examination of the targets and
rods after impact, we inferred the failure mechanisms and speculate on the
properties controlling penetration behavior. The study was aimed at ceramics as a
class of materials and sought a qualitative understanding of penetration
phenomenology to provide the basis for a computational model; therefore,
experiments were performed on a variety of ceramics and details of microstructure
and mechanical properties of the individual materials are not presented.
386 Ceramic Armor Materials by Design
ROD IMPACT EXPERIMENTS
Tungsten-nickel-iron rods (7.70 mm in diameter and 77.0 mm long), having
hemispherical noses, were accelerated in a powder gun to desired velocities in the
0.8 to 1.4 km/s range and caused to impact a steel-encased block of ceramic at
approximately zero degrees obliquity (Fig. 1). Ceramics investigated included
Al2O3, SiC, B4C, and TiB2.
After the impact event, the target assemblies were removed carefully from the
mounting fixture and placed on the floor of the bunker with the impact surface
facing up. A two-component epoxy was poured into the hole on the front cover
plate to fill the crater and run into the cracks in the ceramic. This procedure was
intended to strengthen the fractured ceramic block so that the confining steel
could be removed and the ceramic block could be sectioned without crumbling of
the fractured ceramic. We found in later experiments that impacted targets were
often strong enough to be disassembled, sectioned, and even wafered without
infiltration of epoxy.
Fig. 1. Arrangement for rod impact experiments.
The front surface of the ceramic block and the inside surface of the steel cover
plate showed a starburst pattern of linear markings radiating outward from the
Ceramic Armor Materials by Design 387
impact site. These markings were produced by ceramic and tungsten rod
fragments originating near the leading edge of the penetrator.
All ceramics tested exhibited qualitatively similar cracking patterns; however,
the numbers and sizes of cracks generally differed. The crack patterns on the rear
surfaces of a B4C and a TiB2 specimen impacted at 0.8 km/s are shown in
Fig. 2(a). Two main types of cracks are evident: circular cracks and radial cracks.
Three to four dominant circular cracks were observed in both ceramics; however,
approximately 50% more radial cracks were produced in the B4C (23 cracks) than
in the TiB2 (15 cracks). Fracture damage was heaviest in both materials in a
region directly beneath the impact site.
A diamond saw was used to cut through the crater centers on a plane
containing the impact direction. Cutting the ceramic block was difficult. The
diamond wheel wore out quickly and stopped cutting about halfway through the
block. The wheel needed to be redressed several times before the sectioning was
completed.
Orthogonal views of the craters and crack patterns were obtained on the as-
sawn, unpolished section surfaces. These surfaces for a B4C and a TiB2 specimen
are shown in Fig. 2(b). Crater size was greater in the B4C. Material directly
beneath the crater in the TiB2 specimen remained intact, probably because of the
triaxial compressive stress state and the higher compressive yield strength of
TiB2. At higher velocities resulting in deeper rod penetration, the material in
advance of the tip of the penetrator was crushed to a fine powder. Similar damage
patterns were observed in Al2O3 and SiC targets.
The views in Fig. 2(b) show that the circular cracks in Fig. 2(a) are traces of
cracks that ran outward at an angle from the impact site. The resulting cone
configuration corresponds to the Hertzian cracks observed in the particle impact
experiments reported in the following section and observed by others in ceramics
and glasses under static indentation and particle impact [4, 5]. The lateral cracks
lying roughly parallel to the surface apparently formed after the cone cracks, since
they are discontinuous across the cone cracks. Thus, lateral cracks may be
produced by tensile stress waves reflecting from the specimen boundaries or by
later unloading of the target. Radial cracks are not revealed on cross sections
containing the impact direction.
In several instances tungsten fragments were observed lodged between faces
of Hertzian cone cracks. That these fragments were not moved into the cracks by
the sectioning operation was confirmed by computed tomography results that
showed fragments in cracks in unsectioned specimens. Such observations suggest
that debris emanating from the eroding end of the penetrating rod can have a
significant forward velocity component. These observations support the premise
of Hauver [6], who observed tungsten fragments in advance of the penetrator in
x-radiographs of ceramic blocks during impact by tungsten rods.
388 Ceramic Armor Materials by Design
Fig. 2. Rear surface crack pattern (a) and crack patterns on cross sections (b) in
targets of B4C and TiB2.
Loose fragments produced in the impact experiments were collected for
examination. The size distribution of the collected fragments was determined by
a sieve analysis (placing the recovered fragments on the topmost of a stack of
successively finer screens and vibrating the stack for an hour). We separated
ceramic fragments from penetrator fragments by passing a strong magnet over the
sieved fragments and extracting the slightly magnetic tungsten alloy fragments
from the ceramic debris.
Fig. 3 shows the fragments of SiC retained on screens with various mesh
openings. Fragment shape did not vary substantially with fragment size; aspect
ratios ranged from 1 to about 3. Fracture was predominantly transgranular rather
than intergranular. Differences in fragment size distributions for the four ceramic
Ceramic Armor Materials by Design 389
materials were small over the entire size range. Fig. 4 compares the distributions
for three ceramics in the 2- to 40- m size range.
In higher velocity experiments in which the rod penetrated 60 to 100 mm, the
crushed ceramic material produced at the leading edge of the rod flowed around
and behind the rod, closing the hole made by the rod (Fig. 5). So well
consolidated were these fine fragments that no fragments were loosened during
sectioning and individual fragments were not easily discernable by high
magnification examination. Hardness and scratch tests indicated strengths of the
compacted powder comparable to that of the unimpacted material. In Fig. 5, the
penetrator has stopped just short of the rear confinement plate.
Fig. 3. Fragments from a SiC target retained on screens with different mesh
openings.
390 Ceramic Armor Materials by Design
Fig. 4. Distribution of fragments of B4C, SiC, and TiB2 in the 2- to 40- m size
range.
Fig. 5. Cross section through shot line of 100-mm-thick confined B4C target
impacted at 1.6 km/s showing cracking pattern, compacted ceramic fragments in
cavity produced by penetrator, and embedded penetrator fragments.
The debris from several targets was searched for distal portions of the
penetrator. Distal portions ranging in length from 3-20 mm were found [Fig. 6(a)].
Ceramic Armor Materials by Design 391
Proximal ends had either a mushroom shape or a sharpened-pencil shape. We
speculate that a penetrator tip may alternate shapes between that of a mushroom
and a pencil-point several times during the penetration process. Initially, we
expect plastic deformation of the leading edge to produce a mushroom shape. The
mushroom zone then shears away on roughly a 45° conus, producing a pencil
point. The pencil point then deforms plastically and the tip acquires a mushroom
shape again; this mushroom shape becomes unwieldy and shears to a pencil point.
This alternating shape change continues until penetration ceases. The proximal
end surfaces of all recovered penetrators were faceted and gouged, suggestive of
shear failure. The lateral surfaces were unscored.
Fig, 6. Distal portions of tungsten alloy rods recovered from ballistic experiments
(a) and polished and etched cross section showing deformation of the
microstructure (b).
Polished and etched cross sections on planes containing the rod axis revealed
that the tungsten particles in this sintered alloy retained their original roughly
spherical shape everywhere except near the proximal failure surface. Adjacent to
the failure surface, the tungsten particles were greatly elongated, often to aspect
ratios of 5 or more. Particle distortion decreased with distance from the failure
392 Ceramic Armor Materials by Design
surface, rather gradually (over a distance of about 4 mm) in the mushroomed
region. Thus, the distribution of deformed tungsten particles provides a map of
the plastic strain field in a penetrator.
Tungsten fragments extracted magnetically from the debris ejected from the
impact surface were also examined with a scanning electron microscope. Failure
surfaces and etched cross sections suggested that fragment formation was by
localized shearing of the microstructure, in accord with observations on distal
penetrator ends. Tungsten particle distortion in the fragments, however, decreased
abruptly (usually within about 300 m) from the surface of fragments [Fig. 6(b)].
We computationally simulated the experiment depicted in Fig. 1 using a two-
dimensional, Lagrangian finite difference code. The results provided an estimate
of the distribution and time variation of the stresses and strains produced in the
ceramic target by the impacting long rod before failure occurred, and assisted in
the interpretation of the fractographic observations.
PARTICLE IMPACT EXPERIMENTS
Low-velocity particle impact experiments were performed to study incipient
stages of impact damage. The evolution of fracture damage was established in hot
pressed (HP) silicon nitride by accelerating single solid spheres of tungsten
carbide (WC) or steel onto the polished surfaces of small plate specimens of HP
Si3N4 at a 90°-angle [7]. Particles were accelerated to velocities from 16-368 m/s
by filling the gun chamber with nitrogen gas to various pressures, then suddenly
releasing the nitrogen by rupturing a disk. The diameters of the WC spheres were
1.6 and 2.4 mm; the steel spheres were 2.4 mm in diameter. Impact and rebound
velocities were recorded with photomultipliers. Photomultiplier records also
showed whether particles remained intact or fragmented after impact. The
specimen fracture damage was studied by optical and scanning electron
microscopy of impact surfaces and polished cross sections normal to the impacted
surfaces.
The impact tests caused several kinds of cracks, small craters, and
fragmentation in the target plates and eventually plastic deformation or fracture of
the impacting spheres. Targets sustained no damage at impact velocities below
17 m/s, at which point ring cracks appeared. As impact velocity increased, the
damage progressed to cone cracks, an inelastic impression, radial cracks, lateral
cracks, and median-vent cracks. Ring cracks, as shown in Fig. 7, are
circumferential cracks that extend less than a millimeter beneath the surface. As
the impact velocity increased, more and longer ring cracks formed [Fig. 7(b)].
The ring cracks are similar to the Hertzian ring cracks formed under quasi-static
loading [8]. The surface ring cracks that start approximately normal to the
specimen surface veer outward at various angles up to about 75° from the vertical
to become Hertzian cone cracks [9]. As the velocity increases, additional cone
Ceramic Armor Materials by Design 393
cracks form both inside and outside the existing damage umbrella, and the
innermost cone grows several millimeters in depth.
An inelastic impression and radial cracks [Fig. 7(c)] seemed to form at the
same time in the failure sequence. As the impression deepened with increasing
velocity, the radial cracks grew in both size and number, although only a small
number (8 or 9) of the radial cracks grew to several millimeters [Fig. 7(d)].
Fig. 7. Cracks on the surface of HP Si3N4 caused by impact of 2.4-mm-diameter
tungsten carbide spheres at velocities of (a) 19.5 m/s, (b) 46.2 m/s, (c) 97.7 m/s,
and (d) 159 m/s.
Fig. 8 shows the internal damage and the extent of growth of the various
cracks below the specimen surface. The nucleation and growth sequence of the
ring/cone cracks is illustrated in Figs. 8(a) and 8(b). Under increasingly severe
impacts, cone cracks seemed to cease growing; instead, two new types of cracks
394 Ceramic Armor Materials by Design
were created, as shown in Fig. 8(c). Lateral cracks nucleated internally near the
contact center and ran approximately parallel to, and eventually veered toward, the
impact surface of the specimen. Vertical cracks initiated internally in the region
within the innermost cone crack. These latter penny-shaped cracks are similar to
the median-vent cracks observed by Evans and Wilshaw in quasi-static
indentation experiments on ZnS [8]. Observations with polarized light showed
that a zone of densely microcracked material, approximately spherical in shape,
was formed beneath the contact area. Zinc sulfide impacted by 0.4-mm and 0.8-
mm WC spheres exhibited a similar microcracked zone [10].
Impacting steel spheres, which are softer than tungsten carbide, caused only
ring and cone cracks and introduced little additional damage above 300 m/s, at
which velocity the particle failed by plastic deformation. This limit on the damage
inflicted on the ceramic occurs because the particle cannot exert a pressure on the
ceramic greater than the particle's yield strength. Since the yield strength of the
steel is less than the pressure required for inelastic deformation of the ceramic
surface, higher velocity impacts only result in more deformation of the steel
sphere.
Ceramic Armor Materials by Design 395
Fig. 8. Sectional views of subsurface cracking pattern in HP Si3N4 impacted by
2.4-mm-diameter steel spheres at velocities of 56.4 m/s (a) and 231 m/s (b) and by
a 2.4-mm-diameter tungsten carbide sphere at 231 m/s (c).
396 Ceramic Armor Materials by Design
FAILURE PHENOMENOLOGY OF THE PENETRATION PROCESS
The picture of the penetration process that begins to emerge from these
observations and consideration of the initial stress history is as follows.
Calculations simulations of a tungsten alloy rod impacting a target as in Fig. 1
at 1600 m/s show that at the instant of impact, a shock wave with an amplitude of
several hundred kbars is generated at the impact site. Radial divergence and
plastic flow and fracture in the steel cover plate quickly and drastically reduce the
stress so that the strength of the shock that passes into the ceramic is below the
Hugoniot elastic limit. Thus, the initial shock wave is not expected to condition
the ceramic. A steady-state ramp wave follows the shock, loading the ceramic
material at the tip of the penetrator to a maximum compressive stress of about 50-
60 kbars. The ceramic initially resists the stress in the ramp wave and exerts large
stresses on the tungsten rod, which may deform, fracture, or be deflected.
Ceramics are substantially stronger in compression than in tension, and
consequently, the tensile strength of the ceramic is exceeded at the impact surface
near the rod periphery and tensile fracture begins to occur soon after impact. The
stress fields in the ceramic are initially elastic, and the largest tensile stresses are
in the radial direction (the Boussinesq stress field). Therefore, the cracks that form
(normal to the direction of maximum principal stress) are ring cracks concentric
about the impact site. These cracks are shallow cracks, extending initially only a
millimeter or so beneath the ceramic surface. Upon continued loading, however,
several ring cracks continue to grow and, following the paths normal to the
direction of the principal tensile stress, assume angled trajectories 25-75° outward
from the initial normal-to-the-surface direction. Thus, several large Hertzian cone
cracks extend through the ceramic block, intersect the specimen surfaces, and
cause structural failure of the target.
Up to this point, the stress fields and the fracture response are elastic. But as
the rod continues to advance, the compressive strength is exceeded in material
directly beneath the penetrator. Microcracking occurs in a shallow zone near the
penetrator tip, and the stress field changes in character. The principal tensile
stresses are now in the circumferential direction, and a new type of tensile crack is
invoked six to twelve large radial cracks run outward from the impact site like
spokes from a hub [9]. These cracks intersect the impact surface and may extend
to all specimen boundaries, resulting in strength degradation and eventual
structural failure of the target.
A third crack type, lateral cracks, form beneath the impact surface and
propagate roughly parallel to it, probably during unloading. These cracks intersect
cone cracks and radial cracks, thereby providing the orthogonal surfaces necessary
for fragment formation. Cratering results when these large fragments are liberated
from the vicinity of the impact site.
Ceramic Armor Materials by Design 397
The tensile cone, radial, and lateral cracks do not provide an easy path for
penetrator advance and hence do not assist the penetration process directly. Intact
and laterally confined ceramic remains in the penetrator path despite the presence
of these tensile cracks, and this material must be moved out the way for the
penetrator to advance. This occurs by pulverization of the ceramic material in a
shallow zone immediately ahead of the penetrator and the subsequent flow of this
material laterally and opposite to the impact direction, processes that occur under
large compressive and shear stresses.
Thus, the development of a densely microcracked zone in a ceramic directly
ahead of the impactor is a prerequisite for penetration. Insight into how this zone
forms can be gleaned from the observations of Hagan and coworkers [11, 12] of
damage zones in soda-lime glass and zinc sulfide produced by quasistatic
indentation. These workers observed a curvilinear grid of narrow, fault-like flow
lines beneath indentations, and voids and microcracks at many of the nodes in the
grid. Flow in the polycrystalline ZnS occurred by slip and twinning within the
grains and by grain boundary sliding; voids formed when grain boundary
displacements became large either along the flow lines or at flow line
intersections.
The finely fragmented material at the leading edge of the penetrator wants to
occupy a larger volume (i.e., dilation), but expansion is resisted by the
confinement of the steel encased ceramic block. The resulting increase in pressure
makes fragment flow more difficult and adds to the resistance exerted on the
penetrator. The tensile cracks may assist penetration indirectly by reducing the
level of constraint on the pulverized material, thereby allowing easier flow of the
material out and away from the advancing penetrator, but the main resistance to
penetration is probably coupled to the flow characteristics of highly comminuted
ceramic powder. The cracking pattern in the ceramic target envisioned during the
steady-state phase of the penetration process is depicted in Fig. 9.
As the ceramic particles flow across the leading surface of the penetrator, they
erode the rod, shortening and eventually consuming it as the rod moves through
the ceramic. No scoring or erosion of the sides of the penetrator results from
particles flowing opposite the direction of penetration. Fragments of the penetrator
fretted from the leading surface generally have an initial forward velocity
component and may travel into open cone and radial cracks ahead of the tip of the
penetrator. Other penetrator fragments mix with and flow with the ceramic
powder, becoming part of the front surface ejecta. The eroded tungsten fragments
exhibit greatly elongated grains close to the fragment surfaces, indicative of heavy
localized plastic flow.
398 Ceramic Armor Materials by Design
Fig. 9. Cracking pattern in ceramic targets during the steady-state phase of the
penetration process.
CONCLUSIONS
According to this concept of penetration phenomenology, the properties of a
ceramic that govern penetration resistance include the compressive strength and
hardness, the pulverization characteristics, the frictional flow characteristics of
fine fragments, and fragment abrasiveness. These properties are consistent with
those suggested by Mescall [13, 14].
Initial resistance to penetration is provided by the compressive strength or
hardness of a ceramic. High compressive strength is desirable to deform, fracture,
and deflect an impacting body. Projectiles with low aspect ratios can be defeated
if the strength of the ceramic exceeds the strength of the projectile. High aspect
ratio projectiles such as long rods may suffer heavy deformation and fracture
damage to the proximal end, but the intact distal portion will continue to advance
and penetrate the ceramic. Thus, a high ceramic compressive strength can resist
penetration only to a certain extent.
The stresses exerted by a long impacting rod will eventually pulverize the
ceramic material in a small zone immediately ahead of the leading surface of the
penetrator. As explained in the following paragraph, production of a pulverized
zone is a necessary condition for a penetrator to advance in a confined ceramic
target. Thus, resistance to comminution is desirable for penetration resistance.
Although the specific fracture surface energy for most ceramics is small in
tension, the energy required to produce a unit of failure surface area under large
dynamic compressive and shear forces may be significantly greater. Thus, the
energy absorbed in creating the surface area of the powder may be a significant
ceramic property for penetration resistance.
A penetrator can only advance if the material in its path is pushed ahead of it
or to the side. Because of heavy rear confinement, the crushed ceramic cannot be
pushed ahead and out the rear surface in the way that metallic armor plates fail by
plugging. And if the ceramic is nonporous and snugly confined laterally, the
Ceramic Armor Materials by Design 399
pulverized material cannot be pushed to the side. Indeed, the only recourse is for
the powder to flow opposite the penetration direction along the cavity being
produced by the penetrator. Thus, the pulverized and dilated ceramic must flow
under high pressure, and so the frictional flow property of the comminuted
ceramic should influence penetration resistance. We expect this property to
depend on pressure, strain rate, and size and shape of the fragments, and to be
describable by a Mohr-Coulomb type curve. For thick, highly confined ceramic
blocks, the friction-flow property of ceramic fragments may be the most important
material property for penetration resistance.
Finally, the ability of a ceramic to erode a penetrating rod is a desirable
property for penetration resistance. Whereas erosion may be by gross local plastic
flow of the leading rod surface, ceramic fragments that gouge, score, shear, or
otherwise abrade the rod material may reduce the incoming mass and terminate
the penetration earlier than nonabrasive target materials. Wear and erosion can
occur by a number of mechanisms depending on penetrator and target material,
fragment geometry and size, pressure, temperature, and flow rate. Thus, optimal
erosive behavior might be achievable by matching the abrasive characteristics of a
ceramic material to the threat.
The fractographic observations and the deduced penetration phenomenology
reported here can also be used to identify properties governing the penetration
capability of rods. To be effective as a penetrator, a material should have high
density to produce high stresses in the target; a high yield strength to resist
mushrooming at the leading edge; a high work hardening rate to suppress the
tendency to shear band and fret; a high fracture toughness to resist the propensity
for rod shaft failure; and high abrasion resistance to resist scoring and erosion by
ceramic particles.
In future work, this understanding of penetration phenomenology will be used
to develop tests that measure dynamic shear strength and flow resistance of intact
and fragmented ceramic material under high confining pressure, and to develop
computational models of penetration that can be used to assist in designing
ceramic armor.
ACKNOWLEDGMENTS
Financial support provided by the Defense Advanced Research Projects
Agency and the Army Research Office (Contract DAAL03-88-K-0200), and by
Los Alamos National Laboratory (Contract 9-X69-3295X-1). The authors
gratefully acknowledge the interest and encouragement of Drs. Andrew Crowson,
Eugene Farnum, Francis W. Patten, and William Snowden. Mr. Thomas Cooper
of SRI performed the computational simulations of the rod impact experiments.
400 Ceramic Armor Materials by Design
REFERENCES1F.S. Mascianica, "Ballistic Technology of Lightweight Armor Materials,"
U.S. Army Materials Research Agency, AMRA MS 64-07, Sept 1964 (updated in
1981), AMMRC Report 81-20, Army Materials and Technology Laboratory,
Watertown, MA. 2D. Viechnicki, W. Blumenthal, M. Slavin, C. Tracy and H. Skeele, "Armor
Ceramic-1987," Proceedings of The Third TACOM Armor Coordinating
Conference, February 17-19, 1987, Monterey, CA. 3W. Rafianello, B. Brubaker and R. Hoffman, "Evaluation of a New Low-Cost
Aluminum Nitride Armor," Proceedings of The Fifth TACOM Armor
Coordinating Conference, March 7-9, 1989, Monterey, CA. 4B. Lawn and T.R. Wilshaw, "Review of Indentation Fracture: Principles and
Applications," J. Mater. Sci. 10, 1049-1081 (1975). 5A.G. Evans, "Impact Damage in Ceramics"; p. 302 in Fracture Mechanics of
Ceramics, Vol. 3, Edited by R.C. Bradt, D.P.H. Hasselmann and F.F. Lange,
Plenum Press, New York, 1978. 6G. Hauver, U.S. Army Ballistic Research Lab, personal communication. 7K.C. Dao, D.A. Shockey, L. Seaman, D.R. Curran and D.J. Rowcliffe,
"Particle Impact Damage in Silicon Nitride," SRI Annual Report, Part III, to the
Office of Naval Research, Arlington, VA, N00014-76-C-0657 (1979). 8A.G. Evans and T.R. Wilshaw, "Quasi-Static Solid Particle Damage in Brittle
Solids - I. Observations, Analyses and Implications," Acta Metallurgica 24, 939-
956 (1976). 9A.G. Evans, M.E. Gulden and M. Rosenblatt, "Impact Damage in Brittle
Materials in the Elastic-Plastic Response Régime," Proc. R. Soc. Lond. A 361,
343 (1978). 10D.A. Shockey, K.C. Dao, L. Seaman and D.R. Curran, "Nucleation and
Growth of Cracks in CVD ZnS Under Particle Impact," SRI Annual Report, Part
II, to the Office of Naval Research, Arlington, VA, N00014-76-C-0657 (1979). 11J.T. Hagan, "Shear Deformation Under Pyramidal Indentations in Soda-Lime
Glass," J. Mater. Sci. 15, 1417-1424 (1980). 12S. Van der Zwaag, J.T. Hagan and J.E. Field, "Studies of Contact Damage in
Polycrystalline Zinc Sulphide," J. Mater. Sci. 15, 2965-2972 (1980). 13J. Mescall and C. Tracy, "Improved Modeling of Fracture in Ceramic
Armor," Proceedings of the 1986 Army Science Conference, U.S. Military
Academy, West Point, June 17-20, 1986. 14J. Mescall and V. Weiss, "Materials Behavior Under High Stress and Ultra-
high Loading Rates-Part II," Proceedings of the 29th Sagamore Army Conference,
Army Materials and Mechanics Research Center, Watertown, MA (1984).
Reprinted from International Journal of Impact Engineering, Vol. 9 (3), D.A.
Shockey, A.H. Marchand, S.R. Skaggs, G.E. Cort, M.W. Burkett and R. Parker,
Ceramic Armor Materials by Design 401
"Failure Phenomenology of Confined Ceramic Targets and Impacting Rods,"
pp. 263-275, 1990, with permission from Elsevier Science.
402 Ceramic Armor Materials by Design
MICRO-MECHANISMS OF COMPRESSION FAILURE
Sia Nemat –Nasser and Sai Sarva
Center of Excellence for Advanced Materials
Department of Mechanical and Aerospace Engineering
University of California, San Diego
La Jolla, CA 92093-0416
ABSTRACT
Materials such as rocks, concrete and ceramics fail under different modes
ranging from brittle to plastic failure depending on the deformation conditions.
The under pinning micro-mechanisms of dynamic brittle failure in compression
are examined over a broad range of deformation rates, from quasi-static to strain
rates encountered in ballistic experiments. An overview of recent advances in
novel experimental techniques to study the dynamic behavior of brittle materials
is presented. Recent data on damage initiation and evolution in ceramic armor
materials are considered with a view toward deciphering the essential feature of
failure phenomena. It is observed that under moderate confining pressures and at
moderate deformation rates, brittle failure involves initiation of micro-cracks at
dominant micro-flaws and pre-existing micro-cracks and their subsequent
interactive growth, leading to axial splitting, faulting or a mixture of brittle-
ductile failure. Recent data on SiC is compared to a wing-crack array model,
which describes the influence of microstructure on the dynamic behavior of
ceramics.
Under great confining pressures, common in ballistic impact on the other hand,
classical crack-growth models seem inadequate for representing the actual failure
initiation and evolution. Computational simulations of the early stages of impact
response of ceramic armor show development of stress states involving extremely
high shear stresses within the target ahead of the projectile. This suggests a region
conducive for pulverization. Transmission electron microscopy examination of
recovered Al2O3 powder from a confined sample impact-penetrated by W (X21-
alloy) at high velocity shows extensive twinning with sub-micron spacing.
Corresponding author: [email protected] (858) 534-4914, Fax: (858) 534 2727
Ceramic Armor Materials by Design 403
To the extent authorized under the laws of the United States of America, all copyright interests in this publication are the propertyof The American Ceramic Society. Any duplication, reproduction, or republication of this publication or any part thereof, withoutthe express written consent of The American Ceramic Society or fee paid to the Copyright Clearance Center, is prohibited.
INTRODUCTION AND BACKGROUND
Ceramics, rock and concrete are characterized by brittle failure under
compression. These materials find varied applications based on their mechanical
properties. SiC, Al2O3 and TiB2 find extensive usage in high velocity impact
applications such as multifunctional armor. It is important to understand the
micro-mechanisms of compressive failure so as to help design improved structural
elements. Under quasi-static compressive loading conditions, micro-structural
factors such as mismatches in elastic compliance between adjacent grains and
inherently present processing flaws (e.g. pores and inclusions) create local tensile
stresses. Tensile micro-cracks originate at pre-existing flaws and grow unstably in
the direction of maximum compressive load. Failure occurs by fragmentation
caused by formation and coalescence of these tensile micro-cracks. This mode of
failure is termed as axial splitting. Micro-mechanical models based on pre-
existing flaws, which include frictional and cohesive resistance, have been
presented to describe the failure process. Brace and Bombalakis1 present a sliding
crack model, also termed as wing-crack. The corresponding failure process has
been quantified analytically by Nemat-Nasser and Horii.2 At higher strain rates it
is observed that a many more micro-cracks are nucleated resulting in finer
fragment size. The Hopkinson bar has been modified and extensively used to
study the dynamic behavior of ceramics.3-6
Sarva and Nemat-Nasser7 have
studied the dynamic compressive strength of SiC under uniaxial compression. It
has been observed that the compressive strength of SiC, drastically increases at
strain-rates higher than 1000 s-1
. Some researchers have focussed on the
energetics of nucleation and growth of these microcracks.8,9
It has also been observed that the dynamic properties of brittle materials are
sensitive to confinement.10
In the presence of moderate confining pressure, failure
is inhibited resulting in an improvement in mechanical properties. Failure
eventually occurs by faulting due to a preferential growth of micro-cracks. Horii
and Nemat-Nasser11,12
have suggested that faulting may be a result of unstable
growth of tension cracks at suitable sets of interacting flaws. Chen and
Ravichandran13
have studied the dynamic compressive behavior of a soft ceramic
Macor under confinement.
When confining pressure is extremely large, a transition from brittle failure by
faulting to a ductile response by overall plastic flow takes place. Horii and Nemat-
Nasser12
include possible zones of plastically deformed materials at high shear
stress regions around pre-existing flaws to model the transition from brittle failure
to ductile flow under very high confining pressures. The use of Hopkinson bar to
study the dynamic response under very high confinement pressures is difficult.
However extremely high inertial confinement can be induced in the ceramics by
high velocity impact.
404 Ceramic Armor Materials by Design
Deng and Nemat-Nasser14
have proposed a two dimensional model to simulate
dynamic damage evolution in uni-axial compression. To study the strain rate
effect at moderate strain rates and confinements, Nemat-Nasser and Deng15
have
developed a simple model of an array of interacting wing cracks to describe the
influence of microstructure on behavior of brittle materials at high strain rates. A
brief summary of the wing-crack array model is presented below.
Wing-Crack Array Model
Nemat-Nasser and Deng consider an infinite array of interacting, dynamically
growing wing-cracks.15
The wing-crack array is subjected to a dynamic farfield
compressive load. See Fig.1(a). A bi-axial compressive field is considered to
include lateral confinement. The tensile cracks emanating from the tips of wing-
cracks grow unstably in the direction of the maximum compressive load at limited
speeds. Coalescence of these tensile cracks results in failure. The model is
simplified to an array of collinear cracks as shown in Fig. 1(b). The dynamic
stress intensity factor is calculated by superposition of the solution for a crack
array under pairs of concentrated forces applied at the crack centers and the
solution for a crack array under uniform farfield stresses. The solution to this
collinear crack array displays the influence of the microstructure and the loading
conditions on the dynamic behavior. The microstructure is described by the flaw-
size c and spacing w. See Appendix for the mathematical solution.
Figure 1 (a) and (b). Echelon of wing cracks and collinear crack array model
Ceramic Armor Materials by Design 405
The model helps study the effects of strain-rate, microstructure and lateral
confinement on the compression failure. It predicts that the length scale at which
the material heterogeneities interact with each other leading to micro-cracking and
possible pulverization is dependent on magnitude of compression and strain rate.
At low level of pressure, large defects interact and lead to failure. As the
amplitude of compression is increased, the length scale at which the defects
interact diminishes. The effect of length scale is illustrated in Fig. 2. The model
also predicts that the compressive strength increases with strain-rate and lateral
confinement.
406 Ceramic Armor Materials by Design
Figure 2. The effect of flaw spacing and confinement on the failure stress as
predicted by the Wing-crack array model15
EXPERIMENTAL PROCEDURES
Uniaxial compression
The Hopkinson bar has been extensively used to study the dynamic behavior
of brittle materials at moderate strain-rates and confinement. Ceramics have very
high compressive strengths and low failure strains. Hence, the classical
Hopkinson bar is modified.16,17,18
Since, the behavior is essentially linear elastic,
a pulse shaper, in form of a copper cushion is placed at the front end of the
incident bar to generate a ramp loading and hence maintain uniform strain-rate.
Strain gages are attached to the sample to help measure strains accurately.
Impedance matched platens increase specimen-bar interface area and help reduce
stress concentrations. Previously reported dynamic tests7 of uniaxial compression
of SiC were performed on a 12.7 mm Hopkinson bar. Elastic wave propagation
relations used to calculate the constitutive behavior are valid at the specimen-bar
interface. However, the strain gages that measure the wave pulses are located
Ceramic Armor Materials by Design 407
mid-bar length away from the specimen interface. This shift causes perturbations.
To eliminate the error induced due to these perturbations, the wave pulses are
corrected for dispersion. However, dispersion effects limit the accuracy at high
strain rates (>1500 s-1
) for a 12.7mm Hopkinson bar.
Hence, to attain higher strain rates, a 4.76 mm diameter Hopkinson bar is
used. Further experiments were conducted using cylindrical samples of hot-
pressed SiC. The samples were of 2.03 mm diameter and 3.05 mm length. The
samples were polished parallel to a tolerance of less than 3 m. A striker bar of
38.11 mm length was used to attain a strain rate of nearly 9000 s-1
. Impedance
matched platens were used to prevent bar damage. The W4C platens were
confined in Kovar, to improve their strength. The failure stress was calculated
from the transmitted pulse. Due to the miniature size of the samples, it was not
possible to attach strain-gages onto the sample for accurate measurement of strain.
Hence the strain-rate was inferred from the transmitted pulse and the Young’s
modulus of SiC (470 GPa), assuming that the behavior is linear elastic until
failure. To confirm the validity of the above procedure, experiments were
conducted on a 12.7 mm Hopkinson bar, with strain-gages attached to a SiC
sample. Results indicate that the strain-rate calculations made using the above
technique matched well with the strain-rate measured by the strain gauge attached
to the sample.
For other brittle materials such as concrete, mortar, rock and other relatively
coarse microstructures materials, a 76 mm Hopkinson bar can be used. Strain
rates of nearly 1000 s-1
can be attained.
Moderate confinement
Interference fit technique: An interference-fit, maraging steel sleeve can be used
to achieve a static lateral confinement on the SiC samples. Maraging steel has a
Young’s modulus of 200 GPa and a yield strength of more that 2.34 GPa. This
method consists of fitting a sleeve over the sample, with the sleeve’s inner
diameter smaller than the outer diameter of the SiC sample. Due to the radial
misfit, the sleeve exerts lateral confinement on the SiC sample. Recently,
experiments were conducted to study the compressive strength of SiC under
confinement, using this method. See Fig. 3 for the confinement assembly design.
The radial misfit was about 0.01 mm. The sample was mechanically forced into
the sleeve. Alternatively, they can be shrunk-fit. The confining pressure was
calculated by using the solution for a linear elastic axisymmetric boundary value
problem19
. It was calculated to be approximately 300 MPa.
When the sample-sleeve assembly is axially loaded, it undergoes
compressive strain in the direction of loading. However, it expands in the radial
direction. The Poisson's ratio for the maraging steel is much higher than that of
SiC. Hence, it can be expected that as the axial loading increases, the sleeve
408 Ceramic Armor Materials by Design
expands a larger amount than the SiC sample. This results in the reduction of
lateral confining pressure provided by the sleeve. To counter this, another
maraging steel sleeve, slightly smaller in length was used as an additional sleeve.
The second sleeve was 0.25 mm shorter in length. The inner diameter of this
additional sleeve was chosen to be the same as the outer diameter of the first
sleeve. It helps retain the confining pressure on SiC sample, without inducing any
additional confining pressure. The smaller length of the second sleeve prevents it
from being affected during elastic loading of the sample. Chen and
Ravichandran13
have used a similar technique to laterally confine Macor samples.
The 12.7 mm Hopkinson bar was used to study the dynamic compressive
strength of the confined SiC samples. Since the attachment of strain gauge onto
the sample is not feasible, the strain rate is inferred from the transmitted pulse.
The failure stress was calculated from the transmitted signal. The failure stress
calculation for the confined ceramic is corrected for the inclusion of the metal
sleeve. This adjustment is made by deducting the elastic energy of the metal
sleeve during deformation. The Young’s modulus of maraging steel is 200 GPa.
The sleeve is assumed to be in its elastic regime until the sample fails. This gives
the approximate failure stress for SiC sample.
Figure 3. Sample and confinement design
Pneumatic confinement: For large size samples of materials such as concrete,
rock and polymeric composites, confinement can also be attained pneumatically.20
Fig. 4 shows the 76 mm Hopkinson bar and the pneumatic confinement assembly.
A large diameter pressure vessel, constructed such that it encompasses the entire
sample, provides confinement. The pressure vessel is placed onto the incident and
transmission bar. The proper assembly of pressure vessel is important to ensure
safety and good results. Concrete samples can be tested at confining pressures of a
few hundred MPa and up to a strain-rate of 1000 s-1
, using this system.
Ceramic Armor Materials by Design 409
Figure 4. Pneumatic confining techniques for a 76mm Hopkinson bar20
Figure 5. Tri-axial test configuration21
410 Ceramic Armor Materials by Design
Triaxial tests: The Hopkinson bar can be modified to simultaneously load the
sample in radial and axial directions.21
It consists of larger (27.1 mm) and smaller
(19.1 mm) incident bars and transmission bars as seen in Fig. 5. Incident and
transmission tubes which encompass the smaller incident and transmission bars,
but move independently of them, help load a Teflon sleeve. The Teflon sleeve
surrounds the sample and is restricted by an aluminum sleeve on the outside.
During the test, a large hydrostatic stress is induced in the Teflon sleeve, which in
turn exerts a large radial stress on the sample. The radial stress increases during
the test, as the incident and transmission bars axially load the sample and the
Teflon sleeve. The radial stress is estimated by measuring the hoop strain on the
outer surface of the aluminum sleeve. Strain rates of several thousand s-1
and
radial stresses of several hundred MPa can be attained with this system.
Huge inertial confinement
Experimentally it is very difficult to achieve extremely high confining
pressures or strain rates using the Hopkinson bar. However, large lateral
confinements can be attained by high velocity planar or projectile impact, due to
mass inertia. Nemat-Nasser et al.22
have studied the effect of high velocity impact
by Tungsten X21 alloy on the microscopic failure mechanisms of Al2O3. A 2.54
cm thick Al2O3 tile confined in a steel casing was impact penetrated by W (X21-
alloy) projectile (dia. = 4.8 mm) at about 900 ms-1
. A single stage gas-gun was
used to propel the projectile. During the very initial stages of impact, the sample
is shock loaded and a stress-state similar to that of uni-axial strain exists ahead of
the projectile head. However, as penetration progresses, the emanating stress
waves result in a much more complex state of stress. The ceramic fragments from
the pulverized zone were recovered and examined microscopically by TEM.
RECENT EXPERIMENTAL RESULTS AND FAILURE MODES
Uniaxial compression
As can be seen from Fig. 6, ultra-high strain rate tests indicate that the
uniaxial compressive strength of SiC is approximately 8.5 GPa at a strain rate of
9000 s-1
. Fig. 6 also includes previously reported7 uniaxial compressive strength
data for comparison. Similar to previously reported results, failure occurred by
fragmentation as a result of axial splitting.
Ceramic Armor Materials by Design 411
2
3
4
5
6
7
8
9
10
11
0.00001 0.0001 0.001 0.01 0.1 1 10 100 1000 10000 100000Strain-rate
Com
pressiv
e s
tren
gth
(G
Pa)
Model (confinement = 300 MPa)
Model (unconfined)
Figure 6. Experimental results and comparison to Nemat-Nasser – Deng Model
Moderate confinement
Lateral confinement results in a substantial improvement in the ceramic
strength. The quasi-static failure strength, as measured on an Instron test machine
is 7 GPa. The dynamic strength at a strain rate of 1100 s-1
is 9 GPa. It is seen that
the compressive strength improves by 2 GPa, for a lateral confining pressure of
300 MPa. The strain rate sensitivity of the failure strength is maintained. For the
laterally confined samples, it was observed that the failure is by fault formation
rather than by axial splitting. Fig. 7 indicates the top and bottom view of the
recovered samples. The faults formed were conical in nature. The apex of the
fault cone can be observed in the top view. It runs diagonally across the top face.
The samples were mounted in epoxy resin and ground. The cross section was
observed at regular intervals to examine the failure mode. Fig. 8 is the schematic
of comparison for quasi-static and dynamic failure. It was observed that, for the
same amount of strain, the faults formed for the dynamic test were wider. Also
considerably more microcracking was observed in the dynamic failure as
compared to the quasi-static case.
412 Ceramic Armor Materials by Design
Static Dynamic
Top view
Bottom view
Figure 7. Top and bottom view of the failed samples under moderate confinement
Figure 8. Schematic of faulting for moderately confined samples
Ceramic Armor Materials by Design 413
Large confinement
Extremely high compressive stress, lateral confinement and temperatures are
attained due to shock loading during impact-penetration. The ceramic samples fail
by a combination of failure processes, which include pulverization and
fragmentation due to radial and circumferential cracking. In the pulverized zone,
fragments smaller than the grain size are formed. Transmission electron
microscopy of Al2O3 powder recovered from the pulverized zone indicates
extensive localized plasticity.22
Deformation twins in the sub micron scale were
observed. See Fig. 9 (a). A solitary deformation twin has been isolated and its
electron diffraction pattern studied in Fig. 9 (b). The electron diffraction pattern
displays mirror images of a single pattern super-imposed on each other indicating
that the twin is a ‘reflective-twin’. The characteristic axis of the twin is given by
‘m-m’. Part of the atomic lattice crystal has symmetrically re-oriented itself such
that its structure is a mirror image of the parent matrix lattice. Though Al2O3 has
a HCP crystal structure and a very high degree of symmetry, it is not commonly
known to exhibit twinning under moderate loading conditions. Under extreme
loading conditions, twinning can accommodate extensive plastic deformation with
very little volume change. It is expected that micro-cracks resulting in the
eventual pulverization of the ceramic, accompanies twinning.
Numerical simulations23,24
on DYNA 2D ( a two-dimensional hydrodynamic
finite element code) indicate that during the initial stages of impact, release waves
emanating from the edge of the projectile, produce a zone of high shear stress and
low pressure ahead of the projectile at a distance of the order of its diameter. The
state of stress in this zone lasts only a fraction of microsecond, but is sufficient to
produce pulverization of ceramic.
(a)
414 Ceramic Armor Materials by Design
(b)
Figure 9. Deformation twinning in Al2O3 recovered after impact-penetration by W
(X21-alloy) projectile22
COMPARISON TO NEMAT-NASSER – DENG MODEL
The model has been compared to compressive failure strength data of SiC
for both unconfined and moderate-confinement tests. See Fig. 6. Since the axial
load is substantially larger than the lateral confining pressure, the multi-axial
loading is approximated to bi-axial loading. The model includes the effect of the
microstructure in terms of micro-flaw size and micro-flaw spacing. For
comparison purposes, the model is plotted for pre-existing micro-flaw size of 90
m and flaw spacing of 950 m.
The material constants of SiC are taken as: Young’s Modulus E = 470
GPa; fracture toughness K Ic = 4.5 MPa m ; and the effective Rayleigh wave
speed ms2000Rc-1
. Rayleigh wave speed is the limiting crack velocity. It is
expected that the damage caused by wing-cracks reduces the effective Rayleigh
wave speed of the material. Hence a wave speed, which is approximately third of
an intact material, is chosen. For multi-axial loading, the model is plotted for a
Ceramic Armor Materials by Design 415
confining pressure of 300 MPa. Other parameters defining the wing cracks are
taken as: co-efficient of friction = 0.4; and preferred orientation of micro-flaws
= 72o.
CONCLUSIONS
It is observed that the compressive strength of SiC improves from a quasi-
static strength of 4.2 GPa to 8.5 GPa at a strain rate of 9000 s-1
. Moderate
confining pressures (300 MPa) can be achieved with a ‘two-sleeve interference fit
technique’. The compressive strength of SiC sharply increases under lateral
confinement. A lateral confinement of 300 MPa improves the compressive
strength by about 2 GPa. The strain-rate sensitivity of the compressive strength is
maintained. Under lateral confinement, the failure mode changes from axial
splitting to fault formation. For cylindrical samples, it was observed that conical
faults were formed. Preferential crack growth results in fault formation. An
increased amount of micro-cracking during a dynamic test results in wider faults
during dynamic tests. The Nemat-Nasser – Deng model gives a quantitative
description of the high strain rate behavior of ceramics. It correlates well with
experimental data obtained from the unconfined and moderately confined tests.
Under high amplitude shock compression, the interaction length scale
diminishes to grain size and eventually to nano dimensions. Classical crack-
growth models are no longer applicable under such conditions. In such a regime
even very brittle solids may deform plastically. Thus, comminution may occur
under great confinement through coupling between shear stress, low pressure and
material microstructure. The resulting failure stress will depend on confinement
as well as length scale at which micro-defects interact. TEM results of Al2O3
samples recovered after impact-penetration by W (X21-alloy) show plastic
deformation in form of extensive deformation twinning.
ACKNOWLEDGEMENT
The US Army Research Office supported this project under Contract No.
DAAH04-96-1-0376, to the University of California at San Diego.
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compression,” J. Geophys. Res., 68 3709-3713 (1963)2S. Nemat-Nasser and H. Horii, “Compression induced non-planar crack
extension with application to splitting, exfoliation and rockburst,” J Geo. Phys.
Res., 87 6805-6821 (1982)
416 Ceramic Armor Materials by Design
3J. Lankford, “Temperature-strain rate dependence of compressive strength
and damage mechanisms in aluminum oxide,” J. Mater. Sci., 16 1567-1578
(1981)4J. Lankford, “Mechanisms responsible for strain-rate dependent compressive
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aluminum nitride under uni-axial compression,” J. Mater. Sci., 33 1933-1939
(1998)6C.J. Shih, M.A. Meyers, V.F. Nesterenko and S.J.Chen, “Damage evolution
in dynamic deformation of silicon carbide,” Acta. Mater., 48 2399-2420 (2000) 7S. Sarva and S. Nemat-Nasser, “Dynamic compressive strength of SiC under
uni-axial compression,” Mat. Sci. Engng, A317 140-144 (2001) 8J. Lankford and C. R. Blanchard, “Fragmentation of brittle materials at high
rates of loading”. J. Mater. Sci. 26[11] 3067-3072 (1991) 9D.E. Grady, “Local inertial effects in dynamic fragmentation,” J. Appl. Phys.,
53[1] p825 (1982) 10
J. Lankford, “Dynamic compressive failure of brittle materials under
hydrostatic confinement,” Experimental techniques in the dynamics of
deformable solids, AMD 165 1-11 (1993) 11
H. Horii and S. Nemat-Nasser, “ Compression induced microcrack growth in
brittle solids: axial splitting and shear failure,” J Geophys. Res., 90 [B4] 3105-
3125 (1985) 12
H. Horii and S. Nemat-Nasser, “ Brittle failure in compression: Splitting,
faulting and brittle-ductile transition,” Phil. Trans. R. Soc. Lond. A319 337-374
(1986)13
W. Chen and G. Ravichandran, “Dynamic compressive failure of a glass
ceramic under lateral confinement,” J. Mech. Phys. Solids, 45[8] 1303-1327
(1998)14
H. Deng and S. Nemat-Nasser, “Dynamic damage evolution in brittle
solids”, Mech. Mater., 14 83-103 (1992) 15
S. Nemat-Nasser and H. Deng, “Strain-rate effect on brittle failure in
compression,” Acta Metall. Mater., 42[3] 1013-1024 (1994) 16
S. Nemat-Nasser, J.B. Isaacs and J.E. Starrett, “Hopkinson techniques for
dynamic recovery experiments,” Proc. R. Soc. Lond., A435 371-391 (1991) 17
W.P. Rogers and S. Nemat-Nasser, “Transformation plasticity at high strain
rate in Mg-PSZ” J. Am. Ceram. Soc., 73 136-139 (1990) 18
V. Sharma, S. Nemat-Nasser and K.S. Vecchio, “Dynamic-compression
fatigue of hot pressed silicon nitride,” Expt. Mech. 34[4] 315-323(1994) 19
I. H. Shames and F. A. Cozzarelli, “ Elastic and inelastic stress analysis,”
Taylor and Francis Publishing ltd., p539 (1997)
Ceramic Armor Materials by Design 417
20J. Rome, J.B. Isaacs and S. Nemat-Nasser, “Hopkinson techniques for
dynamic triaxial compression tests,” to appear in Proceedings of Symposium in
honor of I.M. Daniel, Kluwer Academic Publishers, 2002. 21
S. Nemat-Nasser, J.B. Isaacs, and J. Rome, “Triaxial Hopkinson
Techniques,” ASM 8, Mechanical Testing and Evaluation Handbook, 516-518
(2000).22
S. Nemat-Nasser, J.B. Isaacs and B. Kad, unpublished results23
S. Nemat-Nasser and J. Zhang, unpublished results24
S. Nemat-Nasser, S. Sarva, J. B. Isaacs and D.W. Lischer, “Novel ideas in
multi-functional ceramic armor design,” PACRIM IV Conference Proceedings
Maui Nov 4-8, 2001, this volume.
APPENDIX
The Mode I dynamic stress intensity factor at the crack tips in a crack array under
both concentrated and uniform loads is given by
1 21 2 222 2
2
array array array arrayIs IsId Is Is Is
lK k ( l )K k ( l )K , K w tan
w,
(1)
1 2
1
2022 2
2
array array
Is Is
( l l ) lK F cos( ) w sin , K w tan
w w. (1)
As seen from Fig. 1, w is the crack spacing and c is the flaw size. The functions
and , which represent inertia effect during dynamic crack growth in
the stress intensity, are approximately given by
)(1Is lk )(
2Is lk
k (1Is l )
0 75
R
R
c l
c - . l, k (
2Is l )0 5
R
R
c l
c . l
,
(2)
where is the Rayleigh wave speed. where is the driving shear
stress acting on the pre-existing flaw, given byRc 2F c
11 22 11 22 11 22
1 12 2
2 2c( )sin ( )cos , (3)
418 Ceramic Armor Materials by Design
where is the frictional co-efficient, is the cohesive stress describing the wing
crack. The small length is introduced to make the model applicable at = 0.
µ c
0l l
For fracture criterion, it is assumed that the dynamic stress intensity factor
does not exceed a constant fracture toughness, i.e. . The common
growth speed of the compression induced tension cracks is now
arrayIcIdK K
1 2
1 2
1 5 1 75 1 25
1 5 0 75
array arrayIcIs Is
R array arrayIcIs Is
. K . K . K Xl c
K . K . K (4)
where
(5) 1 2
1 2 1 2
2
1 5 1 75 1 25
4 0 5 0 75 0 375
array arrayIcIs Is
array array array arrayIc IcIs Is Is Is
X . K . K . K
K K - K ( . K . K . K ).
The crack length is obtained by integrating the crack tip speed until failure or
complete unloading occurs. To obtain failure stress for a given stress pulse and
material microstructure, the time to failure is calculated from the crack speed and
the length to which the cracks must grow for coalescence to occur. The failure
stress is then defined by the value of the applied axial compressive stress at crack
coalescence.
Ceramic Armor Materials by Design 419
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DAMAGE MITIGATION IN CERAMICS: HISTORICAL DEVELOPMENTS
AND FUTURE DIRECTIONS IN ARMY RESEARCH
D.M. Stepp
U.S. Army Research Office
P.O. Box 12211
Research Triangle Park, NC 27709-2211
ABSTRACT
U.S. Army Research in materials science continues to address the need for
high performance ceramic materials capable of providing superior protection for
the soldier and vehicles in the battlefield; critical to this need is the elucidation of
physical mechanisms that govern the initiation, propagation, and mitigation of
deformation-induced damage. Examples of significant research advances in this
area include such topics as advanced processing, novel characterization tools,
biomimetics, and the altering of comminuted ceramic flow. The U.S. Army’s
objective force, which requires that ceramic armor materials be extremely
lightweight and affordable, serves to further underscore the critical need for
illumination of these governing mechanisms in order to improve ceramic armor
design with robust methods of damage mitigation.
INTRODUCTION
The U.S. Army Research Office has focused on advanced materials since its
earliest days. In 1968, the Metallurgy and Ceramics division summarized the
needs in one of its primary research focus areas as, “…materials having an
appropriate ratio of strength to density, and showing reliability in inhospitable
environments.” Advanced materials were expected to be developed through
investigating the relationships between the phases within a material and its
properties, and between the properties of a material and the principles and
mechanisms that govern them. By 1982, the newly named Materials Science
division identified primary thrust areas in mechanical behavior and synthesis and
processing. Microstructure-property and processing-performance relationships
were expected to yield the future of advanced materials. Today, with the call for
transformation of U.S. Army forces and specific objectives such as the future
Ceramic Armor Materials by Design 421
To the extent authorized under the laws of the United States of America, all copyright interests in this publication are the propertyof The American Ceramic Society. Any duplication, reproduction, or republication of this publication or any part thereof, withoutthe express written consent of The American Ceramic Society or fee paid to the Copyright Clearance Center, is prohibited.
combat system of systems and the objective force warrior, advanced materials
remain at the forefront of Army research.
Although many of the terms and some of the methods have changed, the
underlying need for lightweight materials with superior properties and the
underlying scientific philosophy of developing advanced materials via processing-
structure-property relationships both remain essentially unchanged. These
observations give reason to pause and consider the lessons learned during over
thirty years of ceramics research, that optimal future research directions might be
more clearly identified. Of particular interest is damage mitigation in ceramic
materials, a topic that defines the ultimate figure of merit for armor materials.
Much more than merely penetration resistance, damage mitigation addresses the
need for armor materials to limit the effect of damage on the properties of the
damaged material. As such, proper attention to this capability is essential for a
successful armor materials by design strategy.
BACKGROUND
Processing
The fact that the mechanical properties of ceramics are heavily dependent on
the concentration of agglomerates, inclusions and pores has provided considerable
motivation for exploring optimal and robust methods for processing ceramic
materials. Even early work exploring the mechanical behavior of spinel included
a considerable effort to develop higher purity starting materials and increased
densification [1]. Significant advances in the areas of powder production (purity,
control of particle size distribution, etc.) and processing commercialization
(sintering, slip casting, etc.) have since greatly increased the properties and
availability of a wide variety of ceramics. Although a considerable number of
innovative and promising processing methods have also been developed, they
have generally been limited by cost effectiveness and have not enabled ceramic
armor materials that are vastly superior to those that are available commercially,
particularly in terms of their damage mitigation capabilities.
Mechanical Properties and Characterization
Despite excellent hardness and compressive strengths, the inability to provide
ceramics with substantial toughness has limited their widespread use in many
applications. Although ceramic materials research efforts have demonstrated
substantial strength improvements with additives [2, 3], and microstructural
reinforcements [4, 5], most have not been found to increase the toughness
substantially. Research exploring degradation in the mechanical properties of
ceramics due to such factors as moisture, fatigue loading, and residual stress
concentrations has demonstrated the majority of these mechanisms to ultimately
be governed by the formation and propagation of microcracks [6]. Similarly,
422 Ceramic Armor Materials by Design
thermal shock studies have shown the strength of fractured ceramics is inversely
proportional to the number of cracks that have propagated [7]. It is therefore not
surprising that ceramics incorporating layers, gradients, or confinement, each of
which enhances the ability to inhibit crack formation and impede or deflect crack
propagation, have consistently demonstrated significant property improvements.
In the case of armor, the variations in ballistic performance, inability to
tolerate significant plastic deformation, and propensity for brittle fracture have
severely diminished efforts to develop rugged ceramic armor materials. The
difficulty in achieving this goal is complicated further by the fact that there exists
a significant lack of correlation between quasi-static and high strain-rate
mechanical properties; therefore, even the most substantial improvements in
quasi-static fracture or strength properties tend to provide only marginal, if any,
improvements in ballistic performance. One notable exception to this observation
is confined ceramics, which appear to offer good potential for both structural and
armor applications; however, the cost associated with processing these materials,
particularly with the very large confinement stresses required for ballistic
performance, have been prohibitive to date.
This lack of correlation has provided considerable motivation to develop
experimental tools and characterization techniques to understand the fundamental
mechanisms that govern impact behavior in order to identify and improve upon
salient materials properties and enable the design of superior ceramic armor
materials. The fundamental understanding of the mechanisms by which ceramic
materials deform during impact loading has been advanced tremendously by the
development, modification, and augmentation of such techniques as Taylor
impact, Kolsky and split-Hopkinson bar, plate impact, explosive cylinder, and
spherical cavity expansion. However, depth of penetration (DOP) testing remains
by far the most accepted and predictive method for deriving the ballistic
performance of materials [8]. Nonetheless, accuracy and validity of DOP testing
is difficult to assure, and the technique is often limited and misinterpreted by a
lack of appropriate statistical analysis.
Modeling and Simulation
Due to the difficulties associated with developing robust and valid
experimental tools for predicting ceramic armor material performance, and the
extraordinary costs associated with full-scale armor testing, it is not surprising
that efforts to develop computational predictive models have been numerous. A
wealth of models and simulations have been developed, in many cases
significantly advancing the state-of-the-art in computational theory in order to
address the complexities of penetration behavior in a physically meaningful
manner. At the same time, the desire for increased accuracy has also resulted in
models with required parameters that are so numerous or nonphysical that the
Ceramic Armor Materials by Design 423
results are difficult to interpret. While many simulations have been validated by
comparisons with experimental results, none has emerged with the predictive
capabilities necessary to substantially influence armor design. Further, although
computational models and simulations, particularly when combined with some of
the advanced experimental characterization techniques discussed in the previous
section, enable an enormously detailed analysis of the prevalent
micromechanisms during the penetration process, they have provided no
substantial improvements in damage mitigation capabilities.
CERAMICS IN NATURE
Although no biological organism or system is known to possess substantial
ballistic protection or damage mitigation capabilities, numbers of the robust and
adaptive materials systems found in nature exhibit considerable tolerance for
other forms of deformation damage. For these reasons, the potential exists, albeit
a tremendous challenge, to illumine the mechanisms by which these biological
materials respond and apply them, in a manner appropriate for modern ceramic
materials and high-strain rate response, in order to develop improved ceramic
armor materials capable of mitigating significant damage.
Damage Accumulation in Biological Systems and Ductile Carbides
Analysis of biological ceramic systems has demonstrated an intriguing affinity
for damage accumulation. Both Strombus gigas (Conch) [9] and Haliotis
rufescens (Abalone) [10] shells have been characterized with quasi-static and
dynamic compression and three-point bending tests. In both cases, significant
orientation and strain-rate dependencies were observed. Failure strengths at high
strain-rates (i.e., >103) were measured to be approximately 50% greater than those
measured at quasi-static rates. The materials also exhibited an extraordinary
affinity for damage accumulation prior to failure. Crack deflection, delocalization
of damage, plastic microbuckling (kinking), and viscoplastic deformation were
determined to be the significant mechanisms governing the mechanical response
and enabling the observed damage accumulation.
Recent characterization of bulk Ti3SiC2 has shown a truly remarkable
similarity to the aforementioned biological ceramic systems. These ductile
carbides are characterized by a “layered” unit cell consisting of planer Si layers
connected by TiC octahedra. Mechanical characterization of these materials has
shown significant plastic deformation and the ability to tolerate a considerable
degree of damage. In microstructural analysis, crack deflection, diffuse
microcracking, buckling, delamination, crack deflection, grain push-out, and grain
pull-out were determined to be the significant mechanisms governing the
mechanical response and enabling the damage tolerance. [11, 12]
424 Ceramic Armor Materials by Design
The similarities between the observed deformation mechanisms for these
materials and their ability to tolerate considerable mechanical deformation while
retaining structural integrity are quite surprising. The potential for further
development of layered materials with increased damage accumulation
capabilities, particularly at the scale of the unit cell, appears very strong. Further
work in this area is expected to provide considerable benefit for damage tolerant
structural applications, and possibly for advanced armor materials, if the impact
response of the material can be made to delocalize rapidly enough so as to
mitigate the local stresses and strains within the material.
Lustrin fibers and self-healing
Additional characterization of the abalone shell has led to the discovery of
lustrin, an elastomeric adhesive protein that binds the aragonite plates [13, 14].
Careful analysis of this protein found it to be a modular polymer consisting of
repetitive modular domains. When a strain is applied, these domains enable a
sequential, and reversible, unfolding that provides a “self-healing” response for
the bulk material via sacrificial bonds within the protein, thereby providing
additional fracture resistance and damage mitigation. Although a direct
application of this approach is expected to be of only minimal advantage to an
armor material, the concept is worthy of consideration in order to identify an
appropriate analogue that would enable ceramic materials to substantially mitigate
ballistic damage.
COMMINUTED CERAMIC FLOW
Analysis of penetrated and partially penetrated ceramic materials has led to
the observation of a comminuted zone, also referred to as the Mescall zone, near
the leading edge of the penetrator. Both the resistance to comminution and the
ability of the penetrator to move through the resulting comminuted ceramic
particles have been identified as significant factors governing the ballistic
performance of a ceramic material. In fact, it is primarily on the basis of these
factors that the enhanced performance of confined ceramics has been explained.
In light of these experimental observations, the FRAGBED models were
developed to simulate penetration solely as a combination of fracture,
comminution, compaction, and fragment flow [15, 16]. The model is similar in
principle to atomic dislocation theory, allowing blocky fragments to glide in
discrete increments along fixed material planes, at speeds determined by the local
stresses acting upon them. In addition, fragments are allowed to reduce further in
size by the incorporation of a comminution rate law. This approach to penetration
modeling penetration phenomena demonstrated good agreement with ballistic
testing results, thereby identifying comminuted fragment flow as, potentially, a
Ceramic Armor Materials by Design 425
highly significant mechanism in governing the ballistic performance of ceramic
materials.
More recently, high speed photography and high speed X-ray analysis have
demonstrated that even membrane confinement (e.g., fiber-reinforced packing
tape) of ceramic tiles can substantially alter the shape of the ejecta plume, which
appears to have a significant effect on the ballistic performance of the ceramic
[17]. With minimal front-surface confinement, the ejecta plume has been found
to become less divergent; this focusing of the ejected material appears to have the
effect of further eroding the penetrator, thereby increasing the ballistic
performance of the ceramic. Although this work is preliminary, it provides what
appears to be a novel direction for future ceramic armor research, namely the
mitigation of damage by directing the comminuted ejecta plume against the
penetrator in an optimal manner.
CONCLUSIONS
For more than thirty years, U.S. Army research has investigated and
developed high performance ceramic materials capable of providing superior
armor protection for both the soldier and vehicles. Significant research advances
have been made in low-cost reliable processing and in the characterization and
computational simulation of ceramics during impact. As the Army’s need for
lightweight rugged armor materials becomes increasingly rigorous, efforts to
develop new ceramic materials with significant damage mitigation capabilities
become ever more important.
Despite the numerous advanced processing techniques that have been
developed, cost effectiveness and processing variations continue to severely limit
their applicability to advanced armor materials. Wherever possible, future
processing research efforts should utilize existing commercial processes and focus
on the problems pertaining to the most promising materials solutions (e.g.,
confined ceramics).
While considerable achievements have been made in both experimental and
computational characterizations of ceramic materials, these must be fused to
enable robust predictive armor design tools. One concept of particular interest in
this area is to provide physically-based predictions of the ultimate performance
potential of a wide range of ceramic armor materials and designs, thereby
enabling criteria for eliminating inadequate armor solutions and stimulating
optimum design and processing efforts.
Extraordinary examples of self-healing and damage tolerance capabilities
have been found in biological ceramic systems. Although ductile carbides have
recently been shown to exhibit very similar damage tolerance behavior,
tremendous challenges still remain in elucidating the governing mechanisms in
biological systems and applying them properly; the goal is to obtain a similar
426 Ceramic Armor Materials by Design
level of relative improvement in damage mitigation, but with modern ceramic
armor materials, and during ballistic loading. The potential value of such an
accomplishment would be truly unprecedented.
Finally, recent work has demonstrated the ability to alter the comminuted
ejecta plume and thereby influence the ballistic performance of the ceramic.
Establishing more rigorous control of the ejection process and the potential for
disrupting penetrator performance, including perturbing the path and orientation
of the penetrator, appear to constitute a highly innovative direction for future
ceramic armor research with excellent potential for damage mitigation.
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Controlled Sintering,” in Sintering and Related Phenomena, G.C. Kuczynski,
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Thermomechanically-Processed Polycrystalline Magnesia by Alloying,” in
Structural Ceramics and Design With Brittle Materials, S. Acquaviva and S.
Bortz, eds., Gordon and Breach , New York, 1969. 3J.M. Marder, T.E. Mitchell and A.H. Heuer, “Precipitation From Cubic ZrO2
Solid Solutions,” Acta Metallurgica 31 387 (1983). 4H. Palmour III, “Multiple Slip Processes in Magnesium Aluminate at High
Temperatures,” Proceedings of the British Ceramic Society 6 209-224 (1966). 5J. Hong et al., “Directional Solidification of SiC-B4C Eutectic,” Materials
Research Bulletin 14 775 (1979). 6A Venkateswaran and D.P.H. Hasselman, “Elastic Creep of Stressed Solids
Due to Time-Dependent Changes in Elastic Properties,” Journal of Materials
Science 16 1627-32 (1981). 7B.E. Bertsch, D.R. Larson, and D.P.H. Hassleman, “Effect of Crack Density
on Strength Loss of Polycrystalline Alumina Subjected to Severe Thermal
Shock,” Journal of the American Ceramic Society 57 (6) 235-36 (1974). 8Z. Rosenberg et al., “On the Relation Between the Penetration Resistance eof
Ceramics and Their Dynamic Properties,” Proceedings of the 6th
International
Conference on Mechanical Behavior of Material, ICM 6, 29 July – 2 August
1991, Pergamon Press, 1991. 9R. Menig, M.H. Meyers, M.A. Meyers, and K.S. Vecchio, “Quasi-static and
Dynamic Mechanical Response of Strombus gigas (Conch) Shells,” Materials
Science and Engineering A – Structural Materials Properties Microstructure and
Processing 297 [1-2] 203-211 (2001).
Ceramic Armor Materials by Design 427
10R. Menig, M.H. Meyers, M.A. Meyers, and K.S. Vecchio, “Quasi-static and
Dynamic Mechanical Response of Haliotis rufescens (Abalone) Shells,” Acta
Materialia 48 [9] 2383-2398 (2000). 11
T. El-Raghy, A. Zavaliangos, M.W. Barsoum, and S.R. Kalidindi, “Damage
Mechanisms around Hardness Indentations in Ti3SiC2,” Journal of the American
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M.W. Barsoum and T. El-Raghy, “Room-Temperature Ductile Carbides,”
Metallurgical and Materials Transactions A 30A 363-369 (1999). 13
B.L. Smith et al., “Molecular Mechanistic Origin of the Toughness of
Natural Adhesives, Fibers and Composites,” Letters to Nature 399 [6738] 761-
763.14
X. Shen et al., “Molecular Cloning and Characterization of Lustrin A, A
Matrix Protein From Shell and Pearl Nacre of Haliotis Rufescens,” Journal of
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D.R. Curran, L. Seaman, T. Cooper, and D.A. Shockey, “Micromechanical
Model for Comminution and Granular Flow of Brittle Material under High Strain
Rate Application to Penetration of Ceramic Targets,” International Journal of
Impact Engineering 13 53-83 (1993) 16
J.T. McGinn, R.W. Klopp, and D.A. Shockey, “Deformation and
Comminution of Shock Loaded -Al2O3 in the Mescall Zone of Ceramic Armor,”
Materials Research Society Proceedings 362 61-66 (1995). 17
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428 Ceramic Armor Materials by Design
INDENTATION DAMAGE BEHAVIOR OF ARMOR CERAMICS
Do Kyung Kim and Chul-Seung Lee
Dept. of Materials Science and Engineering
Korea Advanced Institute of Science and Technology
Taejon, Korea
Chang Wook Kim and Soon Nam Chang
Agency for Defense Development
Taejon, Korea.
ABSTRACT
Hertzian indentation technique has been suggested to analyze the damage
response of ceramic materials under the concentrated loading. In the sense that
the impact loading is the specific case of indentation loading, typical armor
ceramics were analyzed by the indentation technique. Experimental indentation
along with numerical calculation was performed to evaluate elastic property, the
yielding stress, and quasi-plastic property. Bonded-interface technique could
provide the observation of subsurface damage pattern after indentation. It is
suggested that quasi-plastic property as well as elastic property is closely related
with the resistance against the ballistic penetration.
INTRODUCTION
Ceramics shows an excellent performance as an armor materials due to its
high hardness and Young’s modulus but low Poisson’s ratio and density. However,
it has not been fully understood which physical, chemical and mechanical
Ceramic Armor Materials by Design 429
To the extent authorized under the laws of the United States of America, all copyright interests in this publication are the propertyof The American Ceramic Society. Any duplication, reproduction, or republication of this publication or any part thereof, withoutthe express written consent of The American Ceramic Society or fee paid to the Copyright Clearance Center, is prohibited.
Figure 1. Hertzian contact of sphere on flat ceramic specimen. Beyond elastic
limit, contact initiates cone fracture (“brittle mode”) or subsurface (“quasiplastic
mode”) [4].
properties influence the armor characteristics of ceramics. Fracture toughness,
strength and hardness data do not provide the enough information to correlate
material properties and armor performance. Even though the dynamic hardness of
ceramics showed some relation with the impact resistance,[1] it is not well known
about the dominant interaction of projectile or shaped charge jet with ceramics.
When a projectile impacts on the ceramics, the stress and the damage
distribution of ceramics are similar to the case of Hertzian indentation, the sphere
indentation on flat ceramics[2, 3]. The sphere indentation on ceramics can provide
the indentation stress-strain relation of ceramics over the wide range of strain. The
compressive stress below the contact area reaches to tens of GPa even with the
normal mechanical testing machine[4, 5], which is comparable to HEL (Hugoniot
Elastic Limit) of the material. It has been suggested that the ballistic efficiency of
armor material is proportional to the average of compressive yield stress and the
HEL of the material[6]. Therefore the characterization of sub-surface microscopic
change of ceramics at a high compressive stress would be helpful to correlate the
materials parameters and the ballistic resistance.
430 Ceramic Armor Materials by Design
In this study the basic mechanical properties of ceramics such as elastic
modulus, Poisson’s ratio, strength, hardness and toughness were characterized.
From the Hertzian indentation method, the microscopic changes were observed
and the indentation stress-strain curves were measured. Yield stress and
strainhardening coefficient were estimated from the numerical analysis based on
the experimental data.
EXPERIMENTAL
In this experiment, well-known eight armor ceramics were evaluated, two
aluminas (AD85, AD90), three silicon carbides (reaction bonded, solid state
sintered and hot pressed), aluminum nitride, boron carbide, and titanium diboride.
The microstructural and mechanical properties are shown in Table I. Specimens
were cut into 3 × 4 × 35 mm and polished. Microstructures were observed using
SEM. Elastic modulus and Poisson’s ratio were measured by pulse-echo method.
Strength was measured by 4-point flexure test and Vickers indentation was used
to measure the hardness and toughness[7].
Figure 1 shows the schematic configuration of Hertzian indentation test. A
spherical ball of radius r is pressed over the flat polished specimen. Beyond a
critical load, either a Hertzian cone crack (“brittle solid”) or a subsurface
deformation zone (“plastic solid”) initiates[4]. At normal load P, the contact
radius a is given by
')'1()1(
16
9,
3
4 223
E
Evvkwhere
E
rkPa (1)
The prime notation denotes the indenter material. The contact radius a defines the
spatial scale of the contact field. The mean contact pressure,
2/ aPpo (2)
defines the intensity of the contact field. From equation (1) and (2), the useful
indentation stress strain relation is defined by
Ceramic Armor Materials by Design 431
r
a
k
Epo
4
3 (3)
Equation (3) means a linear relationship between p0, “indentation stress,” and a/r,
“indentation strain,” leading to a procedure for obtaining basic stress-strain
information. From the contact radius a and load P, the indentation stress and
strain can be experimentally obtained.
Applied load was in the range of 500 and 2000N. Indentation was also made
on “bonded specimen” which was made and polished from two polished
specimens bonded with the glue[8]. Spherical ball was loaded at the exact
position of bonded interface and after detaching the sample using acetone, the
subsurface damage mode was observed. All damage behaviors were observed
using optical microscopy with Nomarski interference after the specimen was
coated with gold.
Finite element computations in this study of elastic-plastic contacts were
carried out using a commercial package (Strand7, G+D Computing, Sydney,
Australia). The configuration is modeled as a sphere of specified radius in
axisymmetric frictionless contact with the flat surface of a half-space, 4 × 4 × 4
mm. Plastic deformation in the test material is governed in our calculations by a
critical shear stress criterion with linear strain-hardening. By imposing a generic
uniaxial compression, stress-strain response of specimen can be described as,
)()(
)(
Y-YEY
YE (4)
where Y is the yielding stress and is the strain-hardening coefficient in the range
0 1 ( = 0, fully plastic; = 1, fully elastic). From this uniaxial
compression, result of numerical indentation stress-strain curves were measured
and compared with the experimental curve. Maximum 50 times of calculations
were iterated to converge the strain-hardening coefficient .
432 Ceramic Armor Materials by Design
Table I. Characterization of microstructure, density and basic mechanical
properties of each specimen
Abbrev. Sample Condition Grain
shape
Grain size
( m)
Density
(g/cm3)
AD85 AD85 Mixed 5 3.44
AD90
Al2O3
AD90 Mixed 2 3.59
AlN AlN Equi-axed 5 3.37
RBSC Reaction
bonded
Bimodal 40,
4
3.08
S-SiC Solid
sintered
Equi-axed 5 3.17
HP-SiC
SiC
Hot
pressed
Equi-axed 2-10 3.22
B4C B4C Hot
pressed
Equi-axed 2-8 2.5
TiB2 TiB2 Hot
pressed
Equi-axed 5-20 4.48
Abbrev. Sample Modulus
E (GPa)
Poisson
ratio,
Strength
(MPa)
Hardness
H (GPa)
Toughness
T
(MPa.m
1/2)
AD85 236 0.230 266 9.2 3.23
AD90
Al2O3
278 0.229 309 12.8 3.19
AlN AlN 327 0.231 288 11.2 2.49
RBSC 394 0.175 440 18.6 3.69
S-SiC 440 0.168 553 29.1 2.46
HP-SiC
SiC
442 0.174 525 19.5 3.75
B4C B4C 456 0.167 390 27.3 3.66
TiB2 TiB2 564 0.081 293 20.6 4.38
Ceramic Armor Materials by Design 433
Figure 2. Half-surface and side views of Hertzian contact damage in (a) AD85,
(b) AD90 (c) AlN (d) RBSC from the WC sphere of radius r = 1.98 mm at P =
1500 N. Reflection optical micrographs of bonded-interface specimens in
Nomarski illumination.
RESULT AND DISCUSSION
Materials Characteristics
Table I shows the result of microscopic characterization and measurement of
basic mechanical properties. TiB2 shows the highest density and B4C shows the
lowest value but the others are similar. In case of Young’s modulus, TiB2 also
shows the highest value. Three kinds of silicon carbides and B4C are the second
highest group. AlN and two kinds of alumina show the lowest modulus, which
means lowest E/d ratio.
Microstructures can be divided as three groups. AlN, S-SiC, HP-SiC, B4C,
and TiB2 have the equi-axed shape and broad size distribution. On the contrary,
RBSC has a bimodal distribution. In case of alumina elongated and equi-axed
grain shape are mixed but the grain size of AD85 is about twice larger than that of
434 Ceramic Armor Materials by Design
AD90. S-SiC and B4C shows the highest hardness. B4C has the highest strength to
weight ratio due to its lowest density.
Figure.2 (continued). Half-surface and side views of Hertzian contact damage in
(e) S-SiC (f) HP-SiC (g) B4C, (h) TiB2 from the WC sphere of radius r =
1.98mm at P = 1500 N. Nomarski optical micrographs of bonded-interface
specimens.
Contact Damage Behavior
Figure 2(a)-(h) compares the section views of the bonded interface including
contact damage of top and side view from the indentation using a WC ball with a
radius of 1.98 mm at 1500 N. Subsurface is observed using bonded specimen
method. Two kinds of opposite behavior are obvious; cone crack mode, and
quasiplastic mode. Ring cracks on surface are connected to subsurface cone
cracks. AD85, AD90, and AlN show the typical damage zone behavior. On the
Ceramic Armor Materials by Design 435
contrary RBSC, S-SiC, and B4C show the typical cone crack damage. HP-SiC and
TiB2 have both kinds of damage characteristics. AD85 – grain size is about twice
larger than that of AD90 – shows larger damage zone than that of AD90. It is
expected that from the analysis of damaged zone the ballistic penetration can be
evaluated.
Figure 3. Side view of damaged zone and FEM-generated stress contours of HP-
SiC. Indentation with WC sphere, radius r = 1.98 mm, load P = 1500 N; (a) side
view of bonded specimen, (b) contour of maximum principal shear stresses with
yielding zone shaded (4.7 GPa at the boundary of shaded zone).
Figure 3 shows the side view of damaged zone with the FEM-generated
contours of maximum shear stress in HP-SiC. The area of stress contour over 4.7
GPa is similar with that of damaged zone. All material constant used for this
calculation are represented in Table II. Yield stress (Y) is defined as the deviation
from the Hertzian elastic limit and can be estimated by experimental data. Strain-
hardening coefficient ( ) is calculated through the iteration of FEM results.
Materials with low show the damaged zone behavior. On the other hand
specimens with high are observed to have the brittle cone cracking behavior.
Figure 4 represents those two opposite characteristics using the indentation
stress-strain relation: (a) in the case of high Y and , (b) in the case of low Y and
. If of a material is zero, it shows fully plastic behavior. And if is 1, it is
considered as a fully elastic material. The result of AD85 is inserted to both (a)
and (b) as a reference date. In figure 4 (a), S-SiC, B4C, and HP-SiC are a group of
showing high Y ( 10 GPa) and high (=0.7-0.8). TiB2 shows a flat graph after
yielding owing to small ( 0.4). In figure 4 (b), RBSC, AD90, AlN, and AD85
show low Y and lower . Each specimen has low yield stress about from 6 to 10
GPa, which is considered to be the limit of elastic regime during impact loading.
436 Ceramic Armor Materials by Design
From the results, indentation stress-stress curve is thought to have some
relationship with impact resistance.
Table II. Elastic and yield parameters for materials used in finite element
modeling
Abbrev. E (GPa) H (GPa) Y (GPa)
AD85 0.230 236 9.2 6.11 0.5
AD90 0.229 278 12.8 7.05 0.6
AlN 0.231 327 11.2 6.58 0.3
RBSC 0.175 394 18.6 6.58 0.6
S-SiC 0.168 440 29.1 9.4 0.7
HP-SiC 0.174 442 19.5 8.93 0.8
B4C 0.167 456 27.3 10.34 0.7
TiB2 0.081 564 20.6 8.46 0.4
WC 0.22 614 19.0 6.00 0.1
Ceramic Armor Materials by Design 437
Figure 4. Hertzian indentation stress-strain curves are plotted for each
specimen. In both (a) and (b), AD85 is inserted as a reference. Data points are
from the experimental measurements. Solid curves are FEM fit as the value of
indicated in Table II.
CONCLUSION
Sphere-indentation technique has been suggested to analyze the damage
response of armor ceramic materials. A special bonded-interface specimen could
provide the observation of subsurface damage pattern after indentation.
Indentation stress-stain curves of each ceramic in elastic-plastic range could be
constructed by the experiments along with the numerical calculations. It is
suggested that quasi-plastic property as well as elastic property is closely related
to the resistance against impact loading.
REFERENCES1D. B. Marshall and A. G. Evans, “Measurement of Dynamic Hardness by
Controlled Sharp-Projectile Impact,” J. Am. Ceram. Soc., 66[8] 580-585 (1983). 2A. G. Evans and T. R. Wilshaw, “Quasi-Static Solid Particle Damage in
Brittle Solids- I. Observations, Analysis and Implications,” Acta Metall., 24, 939-
56 (1976). 3A. G. Evans and T. R. Wilshaw, “Dynamic solid particle damage in brittle
materials: an appraisal,” J. Mater. Sci., 12, 97-166 (1977).
438 Ceramic Armor Materials by Design
4B. R. Lawn, "Indentation of Ceramics With Spheres: A Century After
Hertz," J. Am. Ceram. Soc., 81 [8] 1977-94 (1998). 5B. R. Lawn and T. R. Wilshaw, "Indentation Fracture: Principles and
Applications," J. Mater. Sci., 10 [6] 1049-81 (1975). 6G.R. Anstis, P. Chantikul, B.R. Lawn, and D.B Marshall, “A Critical
Evaluation of Indentation Techniques for Measuring Fracture Toughness: I, Direct
Crack Measurement,” J. Am. Ceram, Soc., 64 [9] 533-538 (1981). 7H. Chai, M. A. S. Kalceff and B. R. Lawn, “Deformation and Fracture of
Mica-Containing glass-Ceramics in Hertzian Contats,” J. Mater. Res., 9 [3] 762-
770 (1994). 8B. R. Lawn, N. P. Padture, H. Cai and F. Guiberteau, "Making Ceramics
"Ductile"," Science, 263 1114-16 (1994). 9A. C. Fischer-Cripps and B. R. Lawn, "Stress Analysis of Contact
Deformation in Quasi-Plastic Ceramics," J. Am. Ceram. Soc., 79 [10] 2609-18
(1996).
Ceramic Armor Materials by Design 439
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PROGRESS IN THE 3-D VISUALIZATION OF INTERIOR BALLISTIC DAMAGE IN ARMOR CERAMICS
Joseph M. Wells*, Nevin L. Rupert, and William H. Green U.S. Army Research Laboratory, Weapons and Materials Research Division Bldg 4600, APG, MD 21005-5069
ABSTRACT The authors present an overview of their research results utilizing X-ray
Computed Tomography (XCT) techniques to nondestructively reveal the internal meso-scale damage morphology within encapsulated armor ceramic targets of TiC, TiB2, Al2O3, and SiC. Examples of the physical damage observed in situ include traditional conical, radial and laminar cracking in impacted samples both with and without penetration. Additional observations reveal instances of outer edge radial cracks, periodic through-thickness laminar cracks, and concentric inplane circular cracking "beach-marks". Examples of asymmetric mixed-type cracking damage isolation and of the spatial distribution of residual tungsten alloy penetrator material are also presented for improved 3-D visualization of complex internal damage conditions. Finally, the authors discuss the premise that this observed meso-scale cracking contributes significantly to the onset conditions for penetration.
INTRODUCTION The physical damage resulting from a high velocity impact of a sub-scale long
rod penetrator with the surface of an armor ceramic is of significant interest to the armor materials community. Even in the case of complete dwell and destruction of the penetrator at the ceramic front surface, i.e. interface defeat, there is considerable damage internal to the ceramic target. With penetration, this damage increases in addition to the growth of a penetration cavity and the deposition of residual penetrator material. Such damage may consist of micro-scale cracking and deformation twinning in a comminuted region immediately under the impact location and of meso-scale cracking in the surrounding elastic ceramic. It is the premise of the authors that the extent and morphology of the latter meso-scale cracking and its resultant structural degradation have a major influence on the cessation of interface dwell and the initiation of penetration.
Ceramic Armor Materials by Design 441
To the extent authorized under the laws of the United States of America, all copyright interests in this publication are the propertyof The American Ceramic Society. Any duplication, reproduction, or republication of this publication or any part thereof, withoutthe express written consent of The American Ceramic Society or fee paid to the Copyright Clearance Center, is prohibited.
Under constraint conditions where an impacted ceramic remains substantiallyintact, it is desirable to characterize the internal damage to understand the nature and location of that damage. Such characterization has been neither effective nor easy to conduct nondestructively. Hence, most prior characterization has been conducted by selective and destructive sectioning and polishing. To better understand the meso-scale (>200µm) details of damage behavior and failure of opaque armor ceramic materials, the authors have applied the nondestructivemethod of X-ray Computed Tomography, XCT. A brief overview of the XCT techniques utilized and several results are presented to demonstrate the innovative and powerful capabilities of this method in the 2-D and 3-D visualization ofinternal damage. Ceramic specimens examined include TiC, TiB2, Al2O3, and SiC. The in situ damage presented occurred predominantly by high velocityballistic impact except for the as-fabricated Al2O3 encapsulated assembly.
IMPACT DAMAGE IN TiCAn evaluation of impact damage in a titanium carbide, TiC, armor ceramic tile
was conducted with details reported elsewhere [1,2]. This sample was impactedwhile confined in a heavy steel encasement that was disassembled prior tosectioning and then a sample half-disk nondestructively scanned for XCTanalysis. A particular feature of this sample is that it was not penetrated, but rather supported extensive dwell or interface defeat of the tungsten alloy penetrator at the impact surface. Nevertheless, appreciable meso-cracking damage was
(a)
(b) (c)
Figure 1. A 3-D Solid Visualization of meso-cracking observed on vertical(b) and horizontal (c) virtual sections of a TiC half-disk sample after interfacedefeat ballistic impact on the front face [1,2].
442 Ceramic Armor Materials by Design
25mm
observed in the interior of this sample. Two quite different types of XCT reconstruction images were prepared to assist in the 3-D visualization of thisdamage. The first is a 3-D solid visualization of the sample half-disk (see figure1a) which has two virtual sections revealing the interior damage on these sections.
The second XCT visualization mode (shown in figure 2) is a reconstructionknown as a 3-D "point cloud" wherein only the selected XCT data related to cracking location, orientation and size are shown. All of the XCT data relating to the opaque ceramic itself has been removed thus allowing the defect cracking to be more easily visualized in isolation. Characteristic features of radial, laminarand conical meso-cracking are clearly observed in the overall asymmetricaldamage condition shown in the front, top and side views of figure 2.
Bifurcated CracksRadial Cracks
Periodic Laminar Cracks
Cone Cracks
Figure 2. X-ray CT Point Cloud virtual representations of 3-D meso-crackingdamage morphology in a TiC ceramic target resulting from ballistic impact. Notethe cracking images are isolated in space without the opacity effects of the TiCmaterial. The outlines of the original sample are superimposed for clarity. [2]
IMPACT DAMAGE IN TiB2
An evaluation of impact damage in titanium diboride, TiB2, armor ceramicdisks was conducted with details reported elsewhere [3,4]. These samples, 72 mm diameter by 25 mm thick, were encapsulated in a titanium alloy, Ti-6Al-4V,welded case and impacted by an L/D=10 tungsten alloy penetrator. Among themore interesting results of this work are:(1) with penetration, the residual
Ceramic Armor Materials by Design 443
penetrator debris was observed in 3-D in a through thickness columnarconfiguration with some dispersion along the side branching cracks (see figures3a & b), (2) considerable meso-cracking damage is observed radially outwardfrom the impact cavity to the outer circumference of the samples (see figure b), (3) radial cracking originating from both the penetration cavity and the outer diameter of the disk was also observed (see figure c) and (4) circumferential orcircular "beach-mark" cracking is also observed in figure 3c.
(a)
(c)
(b)
Figure 3. Virtual 3-D solid sections through the 72 mm dia. x 25 mm thick TiB2
disk (a & b) with penetrator residue and meso-cracking. XCT scan (c) near impactsurface of TiB2 sample showing radial and circular cracks [3,4].
Fabrication Damage in Encapsulated Al2O3
An XCT evaluation of the internal damage in a titanium alloy, Ti-6Al-4V,encapsulated sample with both aluminum oxide, Al2O3, and silicon carbide, SiC,ceramic tiles was conducted in the as-fabricated condition. It is desirable to determine the existence and nature of any initial baseline damage existing prior to ballistic impact.
As shown in figure 4(a), significant damage was revealed in the two left mosttiles in the digital radiograph, DR, taken of the encapsulated sample on edge.Cracking is visible in both of the lower density Al2O3 tiles on the left but not inthe third higher density SiC tile on the right hand side. A XCT scan (see figure
444 Ceramic Armor Materials by Design
4b) taken through the center Al2O3 tile clearly reveals both corner cracks as wellas radial cracking starting in the center and extending outward to the 9 o'clockposition. Such information is useful in the modification of either the targetarchitecture design and/or the fabrication-processing conditions.
(a) (b)
Figure 4. Digital radiograph (a) and X-ray CT scan (b) showing as-fabricatedmeso-cracking damage in Al2O3 tile encapsulated within Ti-6Al-4V sample.
Impact Damage in Encapsulated SiC A preliminary XCT evaluation of the internal damage in a titanium alloy, Ti-
6Al-4V, encapsulated silicon carbide, SiC, ceramic tile sample was conducted in the impacted condition. Partial penetration was experienced with significant lossof ceramic material in the form of a concave cavity. The front face cavity wasfilled with an epoxy resin prior to machining the sample to its reduced size (seefigure 5a). The original sample was too large for XCT with the in-house unit and consequently was machined to 4.75 in square x 0.9 in thick from the originalencapsulation target. As shown in Figure 5b, the damage revealed in the XCT scan at the 13 mm height level (~10 mm from front impact face) in the SiC tileconsists of large asymmetric voids (missing material) and spiral or circular meso-cracking damage rings with several connecting cracks between them. The secondscan image, only 3 mm from the rear face of the tile, is quite different from thefirst. The two scan images are at different depths with a scan thickness of 0.5 mm and a mean separation of 10 mm in the direction of ballistic impact.
Ceramic Armor Materials by Design 445
(a)
(b) (c)
Figure 5. Macro-photograph of epoxy coated Ti-6Al-4V/ SiC impacted sample (a)and two XCT scans: (b) at 13 mm from rear face and (c) at 3 mm from rear faceshowing interior meso-scale damage in 100 cm square SiC tile.
Premise of Critical Damage LevelA general schematic of the types of meso-scale impact damage observed via
XCT in the encapsulated ceramics is represented in figure 6. Damage is definedhere simply as one or more forms of detectable cracking. In simplest terms, threecracking forms observed can be distinguished as conical (or cone cracks), radial(originating from either the center or the outer edge of the target ceramic) andlaminar (parallel to the impact face of the ceramic tile). There is considerableoverlap and merging of these cracking forms in the regions of higher damagedensity and the damage is not necessarily symmetric. The resolution of this XCT technique does not permit the discrimination of the micro-scale damage featuresin either the center comminuted zone or in the surrounding elastic ceramic.
446 Ceramic Armor Materials by Design
With complete dwell and little, if any, penetration, considerable interiordamage can exist in the comminuted region and in the surrounding elastic ceramicregion. For penetration to be prevented by the comminuted region, the structuralsupport (i.e. dynamic confinement pressure) of this region, which is provided and maintained by the surrounding elastic ceramic bulk, must be maintained. Withincreased impact energy of the incoming penetrator, the meso-scale crackingincreases in degree and extent and the structural support provided to the comminuted region decreases. Thus as the meso-scale damage increases, anonunique but "critical" meso-scale damage configuration may be attainedwherein the support to the comminuted region is no longer adequate to maintain its strength to sustain dwell and consequently penetration advances.
Confining Resistance of BulkSurrounding Elastic Ceramic
Ceramic Target Disc LaminarCracks
Cone Cracks
Radial Cracks - ID Radial Cracks - OD
Comminuted Zone
penetrator
Figure 6. Schematic of postulated structural support of the comminuted ceramiczone by surrounding bulk elastic ceramic that allows the comminuted zone toresist penetration
SUMMARY AND CONCLUSIONSX-ray computed tomography has been introduced as a novel and effective
nondestructive methodology to characterize the interior meso-scale damage in penetrator armor ceramics. Damage characterizations described above are post-mortem and are not obtained in real time during the ballistic event. It is thus not possible to directly establish the time sequence of the development of theobserved damage with this technique. It is prudent to conduct baseline DR and XCT scan procedures before, as well as after, impact to capture pre-existingdamage caused by fabrication and handling of target assemblies.
Ceramic Armor Materials by Design 447
Ballistic damage observed with the XCT has included traditional conical, radial and laminar cracking, as well as less-reported spiral or circular beach-mark cracking and outer edge radial and periodic laminar cracking. Resolution of the XCT equipment available for large sample volume sizes of interest prevents the use of XCT for the assessment of micro-scale damage at present. Technology improvements currently under contract are anticipated to provide better resolution by a factor of at least two within present sample size limitations.
It is emphasized that the meso-scale damage observed is postulated to be quite significant in affecting the ballistic performance of the ceramic. Structural support degradation relating to increased meso-cracking surrounding the comminuted zone may be critical to the onset of penetration into the comminuted zone. Future ceramic processing methods introduced to limit or inhibit the meso-scale damage may have a significant benefit in improving the ballistic performance of next generation armor ceramics.
REFERENCES 1.
2.
3.
4.
W.H. Green, and Joseph M. Wells, "Characterization of Impact Damage in Metallic /Nonmetallic Composites Using X-ray Computed Tomography Imaging," pp622-629 in AIP Conference Proceedings 497,1999. J. M. Wells, W.H. Green, and N.L. Rupert, "Nondestructive 3-D Visualization of Ballistic Impact Damage in a TiC Ceramic Target Material," pp159-165 in Proceedings MSMS2001, 2nd Intn'l Conf. on Mechanics of Structures, Materials and Systems, 14-16 February 2001, University of Wollongong, Wollongong, NSW, Australia. W.H. Green, K.J. Doherty, N.L. Rupert, and J.M. Wells, "Damage Assessment in TiB2 Ceramic Armor Targets; Part I - X-ray CT and SEM Analyses," pp130-136 in Proceedings MSMS2001, 2nd Intn'l Conf. On Mechanics of Structures, Materials and Systems, 14-16 February 2001, University of Wollongong, Wollongong, NSW, Australia. N.L. Rupert, W.H. Green, K.J. Doherty, and J.M. Wells, "Damage Assessment in TiB2 Ceramic Armor Targets; Part II - Radial Cracking," pp137-143 in Proceedings MSMS2001, 2nd Intn'l Conf. on Mechanics of Structures, Materials and Systems, 14 - 16 February 2001, University of Wollongong, Wollongong, NSW, Australia.
(Form-CC by Gov Employees - not subject to copyight)
448 Ceramic Armor Materials by Design
Processing and Manufacturing
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AN ASSESSMENT OF LOW COST MANUFACTURING TECHNOLOGY
FOR ADVANCED STRUCTURAL CERAMICS AND ITS IMPACT ON
CERAMIC ARMOR
Richard E. Tressler
Department of Materials Science and Engineering
and the Materials Research Institute
The Pennsylvania State University
118A Steidle Building
University Park, PA 16802
ABSTRACT
The state-of-the-art in manufacturing of advanced structural ceramics,
particularly nonoxides, was recently assessed through visits to several companies
and institutes in Europe and the U.S. Cost of production is a barrier to
widespread application unless the performance is so superior that a cost/benefit
analysis results in favorable economics. The costs of the elements of the
production process are reviewed for specific production processes relevant to
armor production. Target areas for cost reduction for ceramic armor are clear
from this assessment.
INTRODUCTION
The commercialization of new processes and products in the general area of
advanced materials is truly a global enterprise. Technology that is developed
anywhere in the world eventually spreads by licensing, by worldwide marketing,
or by morphing into a variant of the original development. In the areas of
advanced structural ceramics and ceramic matrix composites, the Far East
(particularly Japan), the U.S., and Europe have all been active in developing and
commercializing new technologies. However, the growth in sales of these new
products has not met the expectations of forecasters. Also, the normal price
reductions with increased volume of production have not materialized, further
retarding the realization of mass markets. The military establishment has been
impacted by the high cost of structural ceramic products slowing the widespread
use of ceramics in armor and energy conversion systems. The weakness of market
Ceramic Armor Materials by Design 451
To the extent authorized under the laws of the United States of America, all copyright interests in this publication are the propertyof The American Ceramic Society. Any duplication, reproduction, or republication of this publication or any part thereof, withoutthe express written consent of The American Ceramic Society or fee paid to the Copyright Clearance Center, is prohibited.
pull for these materials and components has had the impact of curtailing
development efforts, particularly in the U.S.
With the slowing of the pace of development and commercialization these
areas in the U.S., the U.S. Army Research Office decided to commission an
assessment of "The State-of-the-Art in Low Cost Manufacturing Technologies for
Advanced Structural Ceramics and Ceramic Matrix Composites" starting in
Europe where some of the leading developments were thought to have occurred in
recent years. Ceramics and ceramic matrix composites have demonstrated great
promise in the production of complex structures with exceptionally high stiffness-
to-weight ratios, chemical stability, impact resistance, and high temperature
capability which are leading to substantial improvements in weight critical and
temperature critical applications. However, the relatively high cost of raw
materials and the complexity of the manufacturing process have created a barrier
to their widespread usage in industrial and defense applications.
This report outlines the findings of this assessment which was conducted in a
two week visit to industrial, academic, government sites in France and Germany,
culminating in a panel discussion at the International Ceramics in Engines
conference held June 19-21, 2000 in Goslar, Germany. Of particular interest for
this assessment were developments in non-oxide ceramics (carbides and nitrides)
and ceramic-matrix composites with both oxide and nonoxide constituents. To
benchmark the findings against comparable industrial activities in the U.S., visits
to five industrial organizations were conducted after the European trip. These
visits focused on industrial firms since the evaluation team was well aware of
research and development efforts underway in academia and government
laboratories through technical meetings and recent NMAB activities (see for
example Reference 1).
The details of the assessment on monolithic structural ceramics are presented
here since this class of materials is of direct interest for ceramic armor. Covalently
bonded, nonoxide ceramics are of special interest because of their low density and
high hardness
OVERVIEW OF MONOLITHIC STRUCTURAL CERAMIC
MANUFACTURING PROCESSES
Most of the manufacturing processes for monolithic structural ceramics start
with powders or powder precursors (sols or gels) and the final densification is
achieved by a sintering or pressure assisted sintering process (hot-pressing or hot
isostatic pressing). Chemical vapor deposition processes have been used to form
polycrystalline ceramic components. Only in the fabrication of certain high purity
semiconductor processing equipment or in preparing coatings are these processes
widely practiced. Some processes involve melt processing (reaction sintering, for
452 Ceramic Armor Materials by Design
example) but generally a powder preform is the starting point for these processes
as well.
Therefore, powder preparation for structural ceramics is the natural starting
point for assessing the manufacturing technologies used for monolithic structural
ceramics.
Powders and Processes for Making Them
The oxides that are important for monolithic structural components are Al2O3
and ZrO2 – alumina because it is still the dominant structural ceramic for a wide
variety of wear, erosion, and impact resistant applications including ceramic
armor, and zirconia because in its toughened form it is an important wear, erosion,
and impact resistant material for use in dies, cutting edges, slitting tools, etc.
Alumina powders for advanced structural ceramics are priced from 1/10 to ¼ of
the price of the least expensive fine grain SiC powder (Carborundum) while
zirconia powders for fine grain advanced ceramics are priced comparably to
submicron SiC powders.
The important nonoxide ceramics are silicon carbide, silicon nitride, boron
carbide, and titanium diboride, and they are presented in order of decreasing
tonnage produced per year. Aluminum nitride is poised to join the top four as it is
becoming more widely used as a thermal management ceramic for electronic
packaging and a erosion/corrosion resistant material for integrated circuit
processing equipment (etchers, for example). Molybdenum disilicide is also
emerging for certain niches such as resistance heating elements. Sheppard has
recently tabulated the various suppliers of fine ceramic powders for advanced
structural ceramics along with the process used to synthesize the powders and the
approximate annual tonnages produced (2).
Silicon carbide is the clear leader in terms of tonnage produced primarily
because there is a large market for abrasive grit. Also, silicon carbide ceramics in
the form of reaction bonded or reaction sintered silicon carbide (where the grains
are bonded with a second phase or phases) have been used for many decades as
specialty refractories for blast furnaces, for metal melting, for porous filters, etc.
Siliconized silicon carbide which is fabricated by infiltrating molten silicon into a
silicon carbide plus carbon preform is widely used for process tubes (for example,
in the semiconductor industry and the metal heat treating industry).
The -silicon carbide powders are made by carbothermal reduction of SiO2
and reaction with carbon. The Acheson furnace which was invented more than
one century ago is still used by all of the major producers of -SiC. Global
production capacity is estimated to be 1 million tons per year (2) with production
levels at about 75% of that figure. The fine sinterable powders are produced
primarily by Norton/Carborundum with Exolon-ESK in second place. H.C. Starck
is a supplier of sinterable powder as is Superior Graphite Co. Both of these latter
Ceramic Armor Materials by Design 453
two companies also market a fine grain -SiC powder but so far the markets for
them have been limited.
Carborundum Structural Ceramics group markets a sinterable premix to a
large number of fabrications. Their production in the calendar year 2000
approached 1 million pounds of premix. They are also a major consumer of this
premix for their Hexoloy family of sintered silicon carbides. Their premix sells
for $11-14/lb depending on the quantity produced.
Silicon carbide powder production is not a roadblock for low cost armor. The
production capacity is larger than the demand worldwide and the price is low
because the fine powder is a side stream of a much larger production of SiC for
abrasive grain and as an additive to metals during melting and refining
Silicon nitride powder is produced by a variety of methods, the most common
being nitridation of silicon powder. Reduction of silicon imide is also used to
make large quantities of sinterable silicon nitride powder. Global production of
silicon nitride is estimated to be several hundred tons. Prices range from $30/kg to
$150/kg depending on particle characteristics (purity, grain size, surface area) and
volume. The major suppliers are Japanese companies (Ubé is a major source) and
European (H.C. Starck).
Boron carbide is produced in a carbothermal reduction process similar to the
Acheson furnace process. The advanced ceramics manufacturers in this country
purchase their powder from ESK. In fact, H.C. Starck resells ESK produced
powder for the advanced ceramics market (including armor). According to
Sheppard's survey there are other suppliers of B4C, which is also used as a
specialty abrasive and as control rod material in nuclear reactors (2). The total
worldwide production is ~100 metric tons. According to ESK the price of B4C
could be comparable to that of SiC if the volume were similar. As it stands now,
hot-pressable powder sells for ~$35-40/kg.
Titanium diboride powders are synthesized by carbothermic reduction of TiO2
and B2O3. There are a number of suppliers (2). The annual production is ~120
metric tons and, thus, the price per kg is high compared to SiC, ranging from $35
to $65/kg.
"Green" Forming of Ceramics
The forming processes employed in the fabrication of parts depends primarily
on the complexity of the shape. Uniaxial dry pressing is used for simple shapes
in volume production because of its reproducibility and use of automation. Free
flowing granules of the powder plus binder, plasticizer, and lubricant are usually
formed by spray drying. Much of the production of armor tiles, wear tiles, pump
seals, etc., use dry pressing for green forming. Dimensional tolerances can be held
accurately; uniformity of composition and, thus, the final microstructure is easy to
achieve, and automatic presses can produce large volumes of parts.
454 Ceramic Armor Materials by Design
Cold isostatic pressing is used for larger diameter cylinders and hollow
shapes. In dry bag pressing the operation is similar to uniaxial dry pressing except
the pressure is applied radially by pressurizing a liquid against a flexible mold
with a rigid shell. At ESK large diameter (~3-4" OD) closed end tubes of Si3N4
were formed this way and green machined to make molten metal handling
equipment.
For complex shapes such as turbine rotors or turbocharger rotors injection
molding is the forming method of choice borrowing from injection molded
plastic technology. However, pressure gel casting, in which a powder-liquids
slurry (slip) is injected into a porous mold under pressure, has been demonstrated
to be amenable to automation and easier to debind by Allied-Signal (Honeywell).
Traditional slip casting is used for large complex shapes with internal
surfaces, particularly in low volume production.
Extrusion is used for constant cross-section products such as tubes and rods.
Carborundum structural ceramics group uses this method to fabricate SiC tubes
for heat exchangers.
Tape casting is used primarily in the electroceramics industry, but is also
used to build up B4C shapes for hot-pressed armor at Ceradyne.
Hot-pressing is clearly an important forming-densification process for
ceramic armor components and for the large more or less flat shapes required for
the semiconductor equipment manufacturers. In most cases a dry pressed or tape
lay-up preform is used to fill the hot-press die cavity. But for single part pressings
it can be filled with free flowing powder. (More on hot-pressing in the next
section.)
Densification
The densification process for the oxide ceramics is straightforward – firing in
air usually with natural gas fired kilns. The densification of the nonoxide ceramics
has been developed in the last 10-15 years so that pressureless sintering of SiC
and Si3N4 based ceramics to 99+% of theoretical density is accomplished
routinely albeit by relatively few companies.
The sintering of silicon nitride ceramics is usually accomplished by liquid
phase sintering. The additives are generally rare earth sesquioxides, often with
Al2O3, and they combine with the SiO2 that is present in all high surface area
powders to form a silicate phase at temperatures of 1750-1900°C. Generally, a
nitrogen overpressure of a couple of atmospheres is required to prevent
dissociation of the Si3N4 and nitrogen loss during sintering.
The sintering of silicon carbide was first commercialized by Carborundum
who used B and C additives which many think altered the surface energetics such
that the material would densify rather than just form necks between adjoining
particles. However, some researchers have also speculated that there is a fugitive
Ceramic Armor Materials by Design 455
eutectic liquid that forms and allows liquid phase sintering. However, the final
product (after ~2000°C firing in inert atmosphere or vacuum) is a 99+ dense –
SiC with some carbon and B4C inclusions. The patents by Carborundum
revolutionized SiC ceramic technology which previously could only be fully
densified by hot-pressing or reaction sintering with excess Si in the final ceramic.
Carborundum and others also experimented with liquid phase sintering of SiC
where Al2O3 and Y2O3 were added, and the firing temperature was above the
eutectic temperature of the oxides. The resulting microstructure contained yttrium
aluminum garnet second phase which resulted in a toughening of the material
although it is not as hard. Carborundum commercialized their product (Hexoloy-
SX) but discontinued it. ESK produces a similar product which is used as wear
plates, primarily in the paper industry.
Hot-pressed silicon nitride and silicon carbide are both being produced by
a number of companies (Carborundum, Ceradyne, Cercom, Kyocera, ESK)
primarily for very high performance applications and where no porosity can be
tolerated (semiconductor processing machinery). Graphite dies are used and
controlled atmosphere or vacuum are required to achieve full density. For the
semiconductor process equipment makers hot-pressing appears to be the only
consistent way to make large parts (18"-20" in diameter) with uniform
microstructure throughout. The other reason for hot-pressing parts for this
industry is that the very high purities required can be achieved since no additives
are needed to achieve densification.
Hot-pressed SiC is also used for armor when very high performance is
required. In general, the hot-pressed products perform better and more
consistently in ballistic tests than the sintered products although it is not clear to
this writer that the latest sintered products have been tested. The Carborundum
enhanced Hexoloy-SA and ESK liquid phase sintered SiC are examples.
The markets for B4C and TiB2 are more limited than SiC and Si3N4. Titanium
diboride is hot-pressed by ESK, primarily for evaporation boats for aluminizing
polymer films and other products. Boron carbide in hot-pressed form is a
lightweight armor material. Tiles of B4C for personnel vests are hot-pressed in
stacks of 20+ at Cercom and 50+ at Ceradyne in large hot-press furnaces that
cycle through the press so that the press is in use essentially full time.
Sintered B4C with metallic additives has been studied for years but no suitable
armor material has evolved from these studies. Carborundum uses a sinter-HIP
(hot isostatic press) process to make complex shapes of B4C and is studying the
process to make B4C helmet armor.
Hot-pressed products in the final machined state cost 2 ½ to 3 times the cost
of sintered products in the final machined state (per Cercom). Much of this
additional cost is the cost of graphite hot-pressing tooling. Machining is a larger
part of the final cost in hot-pressed products. Cheaper hot-pressing is possible
456 Ceramic Armor Materials by Design
(according to Cercom) if a semi-continuous process were developed. For the
armor market none of the manufacturers are willing to invest the capital required
to make the costs lower because the armor contracts are all relatively short term
compared to the time period to recover the additional equipment costs. Thus, to
reduce costs and take much of the manual labor out of the process, much larger
contracts are required or the DOD could fund a Mantech initiative in automated
hot-pressing, or the DOD could develop and own an automated line. No one
armor product volume is large enough to justify the capital costs for a
manufacturer to develop an automated line.
Considerable research is going on now to develop reaction sintering processes
which have the potential to be cheaper and more flexible in terms of incorporating
second phase particles or fibers. Probably the best known process is the Reaction
Bonding of Aluminum Oxide (RBAO) pioneered by Professor Nils Claussen
(now at the University of Hannover). In this process aluminum is incorporated
with aluminum oxide powder and the preform is fired at 1000-1200°C during
which the aluminum melts and oxidizes to form Al2O3. The product can be fully
dense, the forming temperatures are lower than required for Al2O3, and the
process can be net shape through the precise control of volume fractions of Al and
Al2O3.
The process has been applied to Al2O3 with second phase particles such as
SiC; it has been used to make mullite, and researchers around the world are
extending the concept to novel ceramic composites. It has not found widespread
commercial use at this time.
APPLICATIONS
Monolithic Advanced Structural Ceramics
Oxide advanced structural ceramics are used in a wide variety of niches,
primarily where the wear, erosion, and corrosion resistant properties are
important. Coors is probably the largest U.S. producer of oxide structural
ceramics, and they characterize their business as a large number of ~$5M niches.
In general, the wear and erosion resistance of oxide based ceramics is not as good
as SiC and Si3N4 ceramics unless an oxidizing, high temperature environment is
present. In general, SiC and Si3N4 based ceramics have better thermal shock
resistance than oxides with Si3N4 being superior. Silicon carbide ceramics are
better thermal conductors than oxides and most nonoxides with the exception of
AlN. Thus, the oxide ceramics are the low cost choice in many cases while the
nonoxides are the high performance choice. In the case of ceramic armor
materials, alumina is used because of low cost even though it is ~25% more dense
than SiC.
The primary applications for nonoxide ceramics are well-summarized in Table
I. Add to this list the applications identified by ESK (molten metal handling,
Ceramic Armor Materials by Design 457
pump components, evaporation boats) and pursued by Honeywell (nozzle blades
for APUs, aeroengine seals and other components for aeroengines) and one has
most of the niches for application of these materials. Turbocharger rotors have
been in production in Japan but the market is not growing, and diesel engine parts,
particularly the injector parts, are important.
Table I. Primary applications for nonoxide structural ceramics.
Sintered Silicon Carbide
Industrial Seals Heat Exchanger tubes
Auto Seals Semiconductor Equipment Components
Pump Bearings Armor Tile
Wear Tile Kiln Furniture
Silicon Nitride
Bearing Balls Semiconductor Equipment Components
Roller Bearings Wear Parts
Paper Making Equipment Parts Nuclear Seals
Boron Carbide
Nozzles Dressing Sticks
Armor Tile Wear Components
Aluminum Nitride
Electronic Substrates Semiconductor Equipment Components
Source: Saint-Gobain Industrial Ceramics, Structural Ceramics Group, Niagara
Falls, NY
The semiconductor process equipment components are an important market
because the application can tolerate high cost. The hot-pressed, high purity
ceramics are the material of last resort for this class of applications. The total
sales of the major vendors in this market are about $200 million/year which is
probably the largest segment although the segment is composed of many different
parts and configurations.
Meanwhile, investments in processing equipment for the semiconductor
component market is resulting in better processing capability that spills over into
the hot-pressed ceramic armor market. In other words, a commercial market is
subsidizing capital equipment which is used for the ceramic armor market.
Other trends to watch:
The market for seals and bearings is becoming large.
458 Ceramic Armor Materials by Design
Heat exchanger tubes have not developed into a major market, although
there is continuing activity in production of ceramic heat exchangers.
Armor tile is a continuing market, particularly where hot-pressed B4C is
concerned, and components with similar geometry include wear tile for
the paper industry.
Dry-pressed and sintered tiles are in reasonably wide spread production,
and large presses are available or being purchased so that large scale,
semi-automated production of pressureless sintered SiC tiles up to
14"x14" is possible.
There is no semi-automated or semi-continuous hot-pressing of ceramics
at this time, so if hot-pressed armor tile is required by the military, it will
be made in labor intensive, batch process lines. This implies much higher
cost than dry-pressed and sintered material.
Technical and Economic Issues in Manufacturing of
Advanced Structural Ceramics
Advanced ceramic components costs are still too high for widespread
application unless the performance is so superior that a cost/lifetime benefit
analysis results in a favorable economic situation for the ceramics. Elements of
the high cost of production include powder costs, machining costs, and firing
costs. The forming costs are similar to that of powder metals except when hot-
pressing is used, which combines forming and firing.
Powder Costs
For all of the nonoxide ceramics except SiC the powder costs are high. In the
case of SiC the large volume of SiC produced for abrasive, grinding wheels, and
primary metal additives results in lower costs than otherwise would be the case in
view of the relatively small volume of silicon carbide advanced ceramics
produced. The primary cause of high costs of powders is the low volumes
produced. A secondary effect is the stringency placed on powder characteristics –
the higher the purity requirement and the higher the particle size distribution
controls the higher the price (Figure 1). Thus, if the component fabricator can
meet the component specifications with a less stringently controlled powder, costs
can drop. In silicon nitride component production, less expensive powders are
routinely used for wear and erosion components that do not have load bearing
requirements at high temperature.
Ceramic Armor Materials by Design 459
Figure 1. Schematic of powder costs as a function of purity and particle size
control (3).
Machining Costs
Finished machining (fired shapes) must be done with diamond tooling, and all
of the nonoxide ceramics are hard, which means that machining is laborous and
expensive. The key to cutting machining costs is to fire to net-shape so that only
surfaces that require a high finish are machined. There are successes in cutting
machining costs such as pump seal manufacture and ceramic valve machining
where the final machining time was cut to 30 seconds.
Final machining of hot-pressed parts is more expensive than sintered parts
because of the difficulty in holding tolerance during the hot-pressing. Final
machining can contribute 50% to the final component cost for hot-pressed parts.
Firing Costs
Firing of nonoxide ceramics requires temperatures of 1700-2100 C in an inert
atmosphere or vacuum, which is intrinsically more expensive than firing oxides in
air. Continuous kilns have been effective in cutting costs but sufficient volume
must be produced to warrant the continuous operation of kilns. Thus, scale of
production is one of the key factors in firing costs.
Nonuniformity of microstructure across large area parts is difficult to achieve
during sintering due to nonuniformities in green density and binder content and
due to differential rates of heating from edge to center. Additional development
efforts are required to reproducibly produce reliable, sintered armor tiles that will
perform near the level of hot-pressed tiles.
For costs to be cut for hot-pressed parts automated, semi-continuous
processing methods must be developed which requires longer term contracts to
the vendors or direct government investment in automated, semi-continuous lines.
460 Ceramic Armor Materials by Design
Higher Purity/PSD Control
Tota
l Pow
der
Cos
t
The erratic nature of ceramic armor contracts makes it economically unwise for
the vendors to make large capital investments in these types of facilities. Graphite
tooling for hot-pressing is a significant part of the manufacturing cost; there is no
obvious way to substantially reduce these costs.
Reaction based processing of advanced structural ceramics and, particularly,
ceramic matrix composites, holds the promise of reducing costs by reducing firing
temperatures, using cheaper raw materials, shortening processing cycles and
providing near net shape capability. More research and development in this area is
required to commercialize this approach. Examples include reaction sintering of
silicon carbide and liquid silicon infiltration of C/C preforms, and C/SiC preforms
for CMCs. Liquid polymer infiltration as a method to process CMCs warrants
further development.
SUMMARY
Powder costs are high for nonoxide structural ceramics compared to those for
oxide ceramics with SiC being the cheapest of the nonoxides (4-10X the price for
Al2O3). The only way to decrease cost is to increase volume. Machining costs for
these ceramics are on the order of 50% of the total costs except for simple shapes
which can be used with as-fired surfaces. Automated production lines (as in
automobile water pump seals) are necessary for low machining costs.
Firing costs are intrinsically higher for nonoxide ceramics compared to oxides
with temperatures of 1700-1200 C in inert atmospheres or vacuum. Continuous
kilns have been effective in cutting costs but sufficient volume must be produced
to warrant the continuous operation of kilns. Hot-pressing is the process of choice
when uniformity of microstructure across large area parts is required (as in
armor). The cost is 2-3 times the cost of sintered parts. For costs to be cut
automated, semi-continuous hot-pressing methods must be developed which
requires longer term contracts to the vendors or direct investment in such lines by
the customer.
Pressureless sintering of SiC plates has been developed by a few companies to
the state where large area plates can be produced by semi-automated methods
with uniform properties (Weibull moduli approaching 30) and low cost. These
products should be investigated for armor tile and modifications made to attempt
to use these production lines for low cost SiC armor.
ACKNOWLEDGMENT
This work was supported by the U.S. Army Research Office through Batelle
Scientific Services Agreement. The contributions to this assessment by Dr.
Andrew Crowson (ARO) and Dr. James McCauley (ARL) are gratefully
acknowledged.
Ceramic Armor Materials by Design 461
REFERENCES 1Committee on Advanced Fibers for High Temperature Ceramic Composites,
"Ceramic Fibers and Coatings: Advanced Materials for the Twenty-first Century,"
National Materials Advisory Board, National Research Council, NMAB-494,
National Academy Press, Washington, DC, 1998. 2Laurel M. Sheppard, "Global Assessment of High Performance Ceramics for
Armor," report submitted to the Army Research Laboratory, Aberdeen Proving
Ground, MD, 2000. 3D. A. Lathrop, "Non-oxide Powders for Advanced Engineered Ceramics,"
presented at Advanced Ceramics for the New Millenium, March 10-12, 1998,
Atlanta, GA.
462 Ceramic Armor Materials by Design
HIGH-PURITY SUBMICRON -AL2O3 ARMOR CERAMICS
DESIGN, MANUFACTURE, AND BALLISTIC PERFORMANCE
Andreas Krell Elmar Strassburger
Fraunhofer Institut für Keramische Tech- Fraunhofer Institut für Kurzzeit-
nologien und Sinterwerkstoffe (IKTS) dynamik (EMI)
D – 01277 Dresden D – 79588 Efringen-Kirchen
Germany Germany
ABSTRACT
New grades of sintered corundum armor ( -Al2O3) have been designed here to
obtain a ballistic mass efficiency close to SiC and, preferentially, to exhibit a high
optical in-line transmission by associating (i) a small sub- m grain size with (ii)
a very high density and (iii) purity, and (iv) a microstructure that is free of flaws.
Different ceramic technologies like dry (cold isostatic) pressing and casting
approaches (with the option of free shaping) are investigated with respect to these
objectives. Results of ballistic tests give evidence of a strong correlation of pro-
tective efficiency and rising hardness in fine grained sintered Al2O3.
INTRODUCTION
Structural ceramics which associate a high hardness with a low density are suc-
cessfully used as ballistic armor when a high protective power is required at a low
weight. Rankings of the ballistic resistance of different grades of Al2O3, SiC, B4C,
and TiB2 have been established by means of Depths of Penetration (DOP) tests.
However, there is still a lack of fundamental knowledge about the correlation
between the real microstructure of ceramics and their ballistic resistance.
A first systematic study of the influence of materials properties was focused
on alumina ceramics in 1995 and exemplifies the typical difficulties of such in-
vestigations1: the study comprised about twenty commercial Al2O3 ceramics with
different grain size, purity, porosity, and glassy phases, and it was impossible to
analyze suggested influences of individual microstructural parameters (e.g. grain
size) when porosity and glass phase concentration were not constant. Also, the
results showed little correlation between the Hugoniot elastic limit (HEL), the
spalling strength and the ballistic mass efficiency Em. Therefore, only high purity
ceramics with relative densities > 98.5 % should be used in future studies investi-
Ceramic Armor Materials by Design 463
To the extent authorized under the laws of the United States of America, all copyright interests in this publication are the propertyof The American Ceramic Society. Any duplication, reproduction, or republication of this publication or any part thereof, withoutthe express written consent of The American Ceramic Society or fee paid to the Copyright Clearance Center, is prohibited.
gating the influences of grain size or hardness on the ballistic performance.
Previous tests revealed indications that the ballistic resistance of ceramics in-
creases with increasing hardness,2 and it is well known that in polycrystalline ce-
ramics glassy sintering additives reduce the hardness which, on the other hand,
increases with decreasing grain size.3 Starting from these results, it was the ob-
jective of the present work to develop pure Al2O3 ceramics with sub-µm grain
size and to investigate their ballistic performance.
DESIGN OF NEW GRADES OF CORUNDUM ARMOR
To assume a high hardness as a most powerful tool for obtaining a high protec-
tive power seems the more justified as it is commonly agreed that a penetrating
projectile looses a major part of its kinetic energy by deformation and wear inter-
action with the hard armor. On the other hand, it was suggested that on wear there
is a specific hierarchic order of microstructural influences in a way that wear is
more affected by direct influences of the grain size on interface properties (e.g.
reducing pull-out effects by smaller grain sizes) than by the associated hardness.4
In analogy, it may be speculated that a smaller grain size may give some benefit
for an improved ballistic efficiency even when the hardness is not maximum.
Therefore, tests with ultrafine alumina ceramics (grain size < 400 nm) were de-
signed to investigate this issue.
On penetration, the microstructure of the armor collapses within a few micro-
seconds. Therefore, the significance of the usually measured strength of the ce-
ramics is not clear and was addressed here from an empirical point of view.
Among today’s technical Al2O3 ceramics, commercial alumina armors are repre-
sentatives of lower strength grades, often with a 4-point bending strength < 400
MPa. Therefore, the technological efforts of the present investigation were fo-
cused to associate a high macro-hardness close to 20 GPa (at testing load 10 kgf)
with a minimum of flaws in the sintered sub-µm Al2O3 ceramics to enable a high
strength of 500-700 MPa (4-point bending).
On the other hand, these extremely fine grained (sub-µm) corundum micro-
structures that are highly pure and free of defects are also expected to exhibit a
high in-line transmittance of unscattered light (increasing at smaller grain sizes);
the smallest grain size for dense samples was 0.7-0.8 m in these early investiga-
tions with an in-line transmission < 46 % for 1 mm thick disks.5,6
Whereas cubic
materials like spinel can become transparent (“clear”) even with larger grain sizes
as far as the residual porosity is small enough (< 0.05 % requested!), sub-µm
grains are imperative to obtain transparency in hard sintered corundum. Fig. 1
shows the high real in-line transmittance (RIT) obtained now at IKTS Dresden by
eliminating the last residual porosity from the new sub-µm grades of armor.
Contrary to known developments of corundum ceramics that become translu-
cent by a reduced number of grain boundaries per volume (i.e. by grain coarsen-
464 Ceramic Armor Materials by Design
ing associated with a decrease of hardness, protective power, and strength), the
(a) (b)
(c)
Fig. 1. 30-mm discs of sub- m Al2O3 with RIT = 45 % (a, b) and 100x100m2 tile with RIT = 52 %
(all samples 0.8 mm thick; grain size 0.5 m, relative density > 99.9 %; bending strength
650 MPa 4-point - 850 MPa 3-point ). In contrast to translucent armor, transparency is dem-
onstrated here by Fig. 1c and by comparing a polished plate placed (a) immediately on and
(b) in a distance over the printed paper.
sub- m design provides the advantage of combining a greatly improved mechani-
cal performance (cp. data to Fig. 1) with a transition from translucence to a
transparent appearance. Fig. 2 shows the strong increase in the real in-line trans-
mission in a perfect agreement of experimental results and the physical model.7
0
10
20
30
40
50
60
70
80
90
100
0 1 2 3 4 5 6
Experimental data
Calculated model (R. Apetz)
Average grain size (µm)
Real in
-lin
e tra
nsm
issio
n (%
)
7
Ceramic Armor Materials by Design 465
Fig. 2. Influence of grain size on the in-line transmission of sintered Al2O3 (0.8 mm thick samples,
= 640 nm) with a relative density close to 100 %. Physical model7 and measured data.
The real importance of the state of grain boundaries for the ballistic and optical
performance is not finally clear at present. It is, however, worth to note that high
resolution TEM gives evidence that all interfaces of the high-purity alumina ce-
ramics developed here are free of even thinnest amorphous films (Fig. 3).8
Fig. 3. Typical HREM image of a grain boundary in high-purity (>99.9 %) corundum. The
boundary is free of amorphous material or crystalline precipitates.8
PREPARATION OF TILES FOR BALLISTIC TESTS
Ground tiles with a lateral dimension of 100x100 mm2 and with different
thickness (5-15 mm) were prepared from 99.99 % pure Taimicron TM-DAR co-
rundum powder (Boehringer Ingelheim Chemicals, Japan) by
(i) an approach of spray drying and cold isostatic pressing9 or by
(ii) advanced gelcasting8,10
followed by sintering in air; the casting approach offers the additional advantage
of free shaping.10
The samples were prepared without doping additives and by
pressureless sintering if not stated otherwise in Tab. 1.
466 Ceramic Armor Materials by Design
Table 1. Sintered corundum ( -Al2O3) for ballistic investigations. S-samples were pre-
pared by spray drying and cold isostatic pressing, G denotes gelcast materials;
D-samples were provided by Dornier (Friedrichshafen, Germany).
Relative density
(%)
Grain size
( m)
Hardness HV10
(GPa)
Strength4-pt bend
(MPa)
Comments
S-0.3
S-0.5
G-0.6
S-0.7
D-0.9
AD-995
92.5
99.3
100
99.5
98.7
98.8
0.32
0.53
0.57
0.71
0.92
10 - 20
15.0
19.3
20.2
19.1
16.5
12.3
not determined
203 16
644 70
526 55
244 41
350 25
MgO doped; + HIP
Improved non-
aqueous process
Sintering temperatures at 2 hours isothermal hold are about 1420 C for dry
pressed samples and 1260 C after gelcasting to obtain a relative density of 99.5
% (Fig. 4); transparent microstructures require additional hot isostatic pressing.
Fig. 4. Typical microstructure of 99.5 % dense sintered -Al2O3 with 0.54 m grain size.
Table 1 comprises the characteristic data of the samples prepared by different
approaches for ballistic investigations. AD-995 supplied by Coors (Golden, Colo-
rado) was used as a reference which exhibited the highest mass efficiency among
previously tested commercial alumina grades.
BALLISTIC INVESTIGATIONS
Testing set-up and definitions
The different grades of alumina ceramics were tested in a DOP-configuration
(depths of penetration) with a RHA backing (rolled homogeneous steel armor)
of 100 mm thickness and a hardness of HV30 = 3 GPa. The DOP-method was
chosen because it is well established for many years as a method for ranking the
ballistic performance of ceramics,11
and a large body of DOP data exists which
Ceramic Armor Materials by Design 467
can be referred to. Both the surfaces of the ceramic tiles and the steel backing
were ground in order to guarantee reproducible conditions in all experiments.
A tungsten alloy penetrator was selected because this type of projectile is con-
sumed continuously by abrasion. Thus, the scatter of the DOP results is much
smaller than with hard core projectiles which can break or shatter during penetra-
tion. Moreover, a large number of DOP results with that particular projec-
tile/target configuration are available at EMI. In the present investigations, cylin-
drical projectiles with a hemispherical nose (diameter 10 mm, length 32 mm,
mass 44 g) were fired from a 20 mm smoothbore gun by means of plastic sabot
comprising of four petals, an obturator and a steel pusher plate. The impact veloc-
ity was 1250 m/s nominally.
The figure-of-merit for ballistic performance was the ballistic mass efficiency
Em, determined from the residual penetration PR, the penetration into the reference
steel target Pref, the thickness of the ceramic TCer and the densities St, Cer of the
steel and the ceramic. Fig. 5 shows the test configuration and the definition of Em.
RStCerCer
refStm
PT
PE
Fig. 5. DOP configuration and definition of the measured mass efficiency Em
According to Fig. 5, the residual penetration PR observed for a specific ceramic
armor will depend on the thickness Tcer of the ceramic tile and on the densities
Cer and St of ceramic and steel. Usually, a linear decrease of the residual pene-
tration is observed when the ceramic thickness increases resulting in a linear in-
crease of the mass efficiency Em with increasing values of Tcer. From such plots, a
linear extrapolation of Em to a ceramic thickness which would stop the projectile
just at the ceramic-steel interface defines the maximum mass efficiency Em,max as a
characteristic materials parameter.
Experimental Results
The maximum ballistic (protective) mass efficiency Em,max was obtained from
penetration experiments with Al2O3 tiles of different thickness (5-15 mm).
Whereas tests with samples G-0.6 are still in progress, results for the correlation
468 Ceramic Armor Materials by Design
between Em,max and the hardness are given by Fig. 6 for the other grades; a com-
panion paper12
discusses the DOP plots in more detail. Em,max of the commercial
reference AD995 was 2.1 in the projectile/target combination considered here,
typical values for silicon carbide (SiC) are in the range of 3.
Fig. 6 shows significantly higher Em,max values of the fine grained, harder
grades compared with AD995. Plots where Em data for a thickness of 20 mm were
obtained from monolithic tiles or from 10 mm + 10 mm composites yielded
Em,max 2.6 at a hardness of about 18-19 GPa; an even higher value of Em,max =
2.9 was obtained with a 5 mm / 15 mm configuration.
Note that the extremely fine-grained but porous grade S-0.3 with Em,max = 2.3
still exhibits a higher protective power than AD995 - a clear merit of its hardness
which compared to AD995 was increased by the small grain size despite the high
porosity of 7.5 % (cp. Tab. 1).
10 12 14 16 18 20 22 24
Vickers hardness HV10 (GPa)
2.0
2.5
3.0
3.5
no influence of
flaws / strength
Ba
llist
ic m
ass
eff
icie
nc
y E
m,m
ax
AD995
S-0.3
D-0.9S-0.5
S-0.7
G-0.6(expected)
Fig. 6. Influence of the hardness on the maximum ballistic mass efficiency.
The position of the S-0.3 result right on the linear fit of hardness and Em,max in
Fig. 6 excludes any separate influence of the grain size on Em,max beyond the
hardness effect. Hence, Em,max of S-0.3 is lower than that of the coarser but dense
ceramics with grain sizes of 0.5-0.9 µm because here the detrimental effect of the
high residual porosity on the hardness is not balanced by the smaller grain size.
It is important to note that no deviations from the “usual” linear fit in Fig. 6
are observed at constant hardness neither for grades with a low bending strength
(caused by flaws emerging from hard spray-dried granules in S-0.5, cp. Tab. 1)
nor due to different grain sizes: for Em,max it is unimportant whether a high hard-
ness is obtained by a smaller grain size in spite of some residual porosity (S-0.5)
or with a slightly coarser grain size at a higher density (S-0.7).
As to strength effects, however, an influence seems probable for test configu-
rations without the confinement used in the present study (Fig. 5).
Ceramic Armor Materials by Design 469
CONCLUSIONS
The protective power of sintered Al2O3 armor is linearly related to hardness;
there is no separate influence of the grain size or of flaws beyond their impact on
the hardness. The design of new armor ceramics should thus be focused on
- smallest grain sizes significantly below 1 µm,
- high relative density (i.e. minimized residual porosity), and
- high purity.
Additionally, a high strength (e.g. 200 % of today’s commercial Al2O3 armor)
will be beneficial for the general mechanical performance in the technical applica-
tion.
These microstructural conditions are close to the design of new high-strength
transparent sub-µm -Al2O3 (armor) ceramics with a high in-line transmittance.
The new sub-µm alumina grades exhibit significantly higher mass efficiencies
(close to SiC) than commercial corundum armor tested under the same conditions.
Free shaping of these armor components is enabled by new casting approaches.10
ACKNOWLEDGEMENTS IKTS Dresden gratefully acknowledges the cooperation with Dr. R. Apetz and Dr. M.
van Bruggen at Philips NatLab (Eindhoven, NL) within the STARELIGHT project
funded by the European Commission (“Growth” program, contract G5RD-CT-1999-
00088).
REFERENCES 1B. James, “The influence of the material properties of alumina on ballistic
performance,” pp. 3-9 in Proceedings of the 15th
International Symposium on Bal-
listics (Jerusalem/Israel, 1995 published by the Organizing Committee).2I. Faber, K. Seifert and L.W. Meyer, “Correlation between the mechanical
data of ceramics and their protective power against impact loading” (in German),
Final Report EB 6/95 (part 3), Technical University Chemnitz-Zwickau, Depart-
ment of Engineering Materials, 1995. 3A. Krell and P. Blank, “Grain Size Dependence of Hardness in Dense Submi-
crometer Alumina,“ J. Am. Ceram. Soc. 78 4 1118-20 (1995).4A. Krell, “Improved hardness and hierarchic influences on wear in submicron
sintered alumina,“ Mater. Sci. Eng. A 209 1-2 156-63 (1996). 5K. Hayashi, O. Kobayashi, S. Toyoda and K. Morinaga, “Transmission optical
properties of polycrystalline alumina with submicron grains,“ Materials Transac-
tions (JIM) 32 11 1024-29 (1991). 6H.Mizuta, K. Oda, Y. Shibasaki, M. Maeda, M. Machida and K. Ohshima,
“Preparation of high-strength and translucent alumina by hot isostatic pressing,“
J. Am. Ceram. So.) 75 2 469-73 (1992).
470 Ceramic Armor Materials by Design
7R. Apetz and M. van Bruggen, “Transparent Alumina: a Light Scattering
Model,” submitted to J. Am. Ceram. Soc.8A. Krell, E. Pippel, J. Woltersdorf and W. Burger, “Subcritical crack growth
in sub-µm Al2O3,” J.Europ. Ceram. Soc. (in press - 2002). 9E. Strassburger, H. Senf, H. Rothenhäusler, B. Lexow and A. Krell, “Influ-
ence of grain size and microstructure on the ballistic resistance of alumina,” pp.
1216-23 in Proceedings of the 18th
International Symposium on Ballistics (San
Antonio/TX, 1999), Technomic Publishing Co., Lancaster/PA, 1999. 10
A. Krell, „High-strength Al2O3 joint prostheses of complex shape,“
http://www.
ikts.fhg.de/business/strukturkeramik/basiswerkstoffe/oxidkeramik/al2o3_bio_eng.
html.11
Z. Rosenberg, S. Bless, Y. Yeshurun and K. Okajina,“A new definition of
ballistic efficiency of brittle materials based on the use of thick backing plates”,
pp. 491-98 in Impact Loading and Dynamic Behaviour of Materials (Proc. Impact
87 Conf., Bremen, Germany, 1987), DGM Informationsgesellschaft, Oberursel,
1988.12
E. Strassburger , A. Krell, B. Lexow, „Ceramic Armor with Submicron Alu-
mina against AP Projectiles,“ pp. Xx-xx in Proceedings of PAC RIM IV, Ceramic
Armor Materials by Design (Wailea, Maui, Hawaii, 2001).
Ceramic Armor Materials by Design 471
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SOLID FREEFORM FABRICATION OF ADVANCED ARMOR CONCEPTS:
OPPORTUNITIES FOR DESIGN AND MANUFACTURE
R.C. McCuiston, S.C. Danforth, M.J. Matthewson, and D.E. Niesz
Malcolm G. McLaren Center for Ceramic Research
Rutgers, The State University of New Jersey
607 Taylor Road
Piscataway, NJ 08854
ABSTRACT
There is tremendous interest in advanced armor concepts. Fortunately, there
are novel manufacturing methods available for such systems, referred to as Solid
Freeform Fabrication (SFF), or Layered Manufacturing (LM). These methods are
generally free of the normal constraints imposed by traditional manufacturing.
Designers now have the capacity to optimize design for performance in ways
never before possible. Using SFF or LM technologies, one can manufacture
components out of several different materials to achieve multi-functionality. This
is accomplished by controlling the spatial distribution of materials by a computer
driven material deposition system. A brief review of our SFF method, Fused
Deposition of Ceramics (FDC) will be given along with possible novel armor
design concepts.
INTRODUCTION
In what seems to be an ever-shrinking world, the need to travel around it on a
moments notice is all the more important. This is especially true in times of war
and conflict. One of the expressed goals of the United States Army is to be able to
deploy to anywhere in the world from a multitude of dispersed sites in a matter of
days. Unfortunately with the increasing lethality of today’s weapons, designers of
vehicle platforms have had to compensate by adding increasingly thicker armor,
typically dual hardness steel. As a result air transportation of these increasingly
heavier vehicles is very difficult. The added weight also increases fuel
consumption, decreases maneuverability and tests the limits of portable bridges.
[1]
To help realize the goals of an easily deployable force, research has been
conducted on novel armor concepts such as confined ceramic tiles [2], confined
Ceramic Armor Materials by Design 473
To the extent authorized under the laws of the United States of America, all copyright interests in this publication are the propertyof The American Ceramic Society. Any duplication, reproduction, or republication of this publication or any part thereof, withoutthe express written consent of The American Ceramic Society or fee paid to the Copyright Clearance Center, is prohibited.
multi-layered ceramic metal systems [3,4] as well as unconfined multi-layered
ceramic metal composites. [5] In the first two examples, it has been shown with
depth of penetration tests (DOP), that an external confinement, in which either a
lateral, or hydrostatic pressure is applied via a metal phase, provides improved
ballistic performance over unconfined armor systems. It is known that a highly
cracked region of ceramic directly in front of a projectile is needed in order for a
penetration event to initiate. [4] This highly cracked region will develop into a
comminuted zone and the projectile will penetrate by forcing comminuted
fragments to flow around the advancing projectile and thus becoming ejected
from the impact site. Under external confinement however, this flow, and
subsequent ejection of comminuted fragments is hindered, allowing the
comminution zone to aide in projectile erosion. Unfortunately, it is impractical to
use these externally confined armor systems on a large scale. The edges of these
armor systems are pure metal, creating unprotected areas when a single layer of
tiles is applied.
In the past, metal matrix composites (MMC) have been shown to have
improved ballistic performance, this being attributed to dynamic work hardening.
[1] The work hardening of the metal was limited due to microstructural damage
created by shockwave interactions. It was thought that creating a multi-layered
ceramic metal composite, where each layer would contain different percentages of
ceramic, might further improve ballistic performance. [5] Further research is still
required, however, as many fundamental questions, such as what layering design
and what size scale is critical for optimal shockwave mitigation.
It is increasingly apparent that new concepts in armor design, as well as new
methods to rapidly create them to allow for multiple design iterations is needed.
This paper will discuss several new armor design concepts, as well as SFF
manufacturing methods for them. Some preliminary results are presented which
show feasibility for fabricating these new concepts by FDC.
ARMOR DESIGN CONCEPTS
If a level of confinement is to be provided to the ceramic phase, without using
an external method, some form of internal confinement via a reinforcing phase
must be applied. Infiltrating a porous ceramic perform with a molten metal,
creating a metal matrix composite, might provide some degree of internal
confinement. Evans et. al [6] has shown that periodic metal structures, when
designed properly, have improved properties over that of stochastic metal
structures. It stands to reason then, that a purposely-designed internal
reinforcement phase should provide improved properties and thus performance,
over that of randomly created reinforcement phase.
Work by Rödel et. al [7] and Claussen et. al [8] on alumina / aluminum
composites has shown that alumina reinforced with fibers of aluminum had the
474 Ceramic Armor Materials by Design
same fracture toughness as alumina performs infiltrated with molten aluminum.
However the fiber reinforced alumina only required 13 volume % aluminum
whereas the infiltrated alumina contained 25 volume % aluminum. The fiber
reinforced system obviously allows for a much better improvement in properties,
while using less reinforcing phase.
Figure 1 is a schematic of a fiber reinforced armor composite concept. It
should be noted that to allow for easy visualization, the impact face has been
placed towards the bottom. The light gray regions are the ceramic phase and the
dark gray regions are the metal fiber reinforcement. Design flexibility to enable
testing of multiple designs is quite large. The diameter and placement of the metal
fibers, as well as their volume fraction can all be tailored for optimal properties,
when using FDC.
Figure 1.) Schematic of an internally reinforced ceramic metal armor
composite.
Figure 2 is a schematic of another possible armor composite that would utilize
shockwave mitigation as a means of improved performance. [9] The light gray
region is a continuous phase, while the darker spheres are a discontinuous phase.
Chin et. al [5] have stated that macroscopic interfaces in layered armor
composites are extremely important. It is thought that these interfaces will play a
role in controlling the reflection and refraction of shock waves during impact
events and could, if designed properly, be used to essentially steer the stress
waves and improve performance. Figure 2 provides many size scales of
interfaces, to enable the control of various frequency shock waves, and is easily
tailorable by changing the size of the spheres, their stacking order, and volume %.
Ceramic Armor Materials by Design 475
Figure 2.) Schematic of an armor composite showing many scales of
macroscopic interfaces.
FUSED DEPOSITION OF CERAMICS
To fabricate and test multiple design iterations of the armor composites shown
in Figures 1 and 2, both rapidly and accurately, one would ideally want to use
some form of a solid freeform fabrication (SFF) technique. There are several SFF
techniques capable of producing functional ceramic components, among them are,
Stereolithography [10], 3-Dimensional Printing [11], Selective Laser Sintering
[12], Robocasting [13], and Fused Deposition of Ceramics (FDC) [14].
It has been shown through prior work with ISR-Si3N4, that FDC is capable of
producing functional components. [14] An average four point bend strength of
908 MPa and a chevron notch fracture toughness of 8.5MPam1/2
were measured
on FDC Si3N4 bars, which is comparable to commercial Si3N4. Moreover, the
bend strength and fracture toughness were not statistically different when
measured parallel and perpendicular to the build layers, indicating that FDC
produces nearly homogenous parts. Due to its extrusion based technology
however, FDC can be used to introduce crystallographic texture. By adding -
Si3N4 seeds to filament feedstock, preferred grain orientation was observed in
FDC Si3N4. [15] Figure 3 is a schematic of the FDC process.
476 Ceramic Armor Materials by Design
Figure 3.) Diagram of the Fused Deposition of Ceramics process.
The FDC process works by extruding a ceramic loaded thermoplastic
filament, through a fine nozzle. The roads are laid down in the x-y plane in a
controlled fashion until a single layer is completed. A z-stage then indexes down
one layer thickness and another layer is built on top of the previous layer. A
complete description of the FDC process is given elsewhere. [16-18]
Subsequently, after FDC part fabrication, binder removal and sintering steps are
performed.
An advantage to using FDC as a fabrication technique for new armor design
concepts is that it can spatially distribute material in the x, y and z planes. With a
multiple extrusion head FDC system, one can also spatially distribute multiple
materials in all three planes, lending another tool to the design of these new
concepts.
MODEL MATERIAL SYSTEM
Research has been initiated to study the effect that reinforcing metal fibers
have on the impact performance of ceramic metal armor systems. A model system
of alumina and copper has been selected to allow for relatively easy fabrication,
and thus rapid design iteration. FDC has been used to fabricate several alumina
scaffolds containing designed channels for molten metal infiltration. These
scaffolds will then be spontaneously infiltrated with a wetting copper-oxygen
alloy to create a confining fiber phase. [19,20].
Ceramic Armor Materials by Design 477
Initial work has been done using FDC filaments containing 55 volume %
Alcoa 152-SG alumina. Figure 4 is an SEM image showing the cross section of a
sintered alumina scaffold produced by FDC. This scaffold was designed with a
volume fraction gradient through the thickness and it is apparent that the channel
volume is relatively uniform in each layer.
Further work was done to show that infiltration of a sintered alumina scaffold
was feasible. Figure 5 is a light optical image of a sintered alumina scaffold
spontaneously infiltrated with copper. This sample was produced by filling the
sintered scaffold with copper powder, and then melting it under static air. The
copper alloyed with oxygen in the air and subsequently wet and infiltrated the
scaffold. While this is by no means an ideal method of infiltration, it does show
that ceramic metal reinforced armor composites can be fabricated using a
combination of FDC and spontaneous infiltration.
Figure 4.) SEM image showing a cross section of a sintered alumina scaffold.
478 Ceramic Armor Materials by Design
Figure 5.) Light optical image showing an infiltrated alumina scaffold. The
dark ovals are the alumina scaffold, while the lighter phase in between is the
copper-oxygen alloy.
SUMMARY
New armor design concepts are needed to help solve the externally confined
ceramic armor problem as well as improve upon armor performance by
shockwave mitigation. It is thought that providing a purposely designed, internal
reinforcement phase might provide a degree of internal confinement. It is further
thought that tailoring of macroscopic interfaces in armor composites to mitigate
stress waves is another approach. The use of FDC, along with metal infiltration
has been shown to be feasible way to rapidly design iterate and fabricate novel
internally reinforced ceramic armor composites.
ACKNOWLEDGEMENTS
The authors would like to thank the U.S Army Research Laboratory for
funding under cooperative agreement number DAAD19-01-2-0004, as well the
CCMC for additional support. We would also like the thank Dr. McCauley, Dr.
Adams, and Dr. Chin of the ARL for technical input.
REFERENCES1E.S.C. Chin, “Army focused research team on functionally graded armor
composites,” Materials Science and Engineering A, 259 [2] 155-61 (1999).
Ceramic Armor Materials by Design 479
2C.E. Anderson Jr. and S.A. Royal-Timmons, “Ballistic performance of
confined 99.5%-Al2O3 ceramic tiles,” International Journal of Impact
Engineering, 19 [8] 703-13 (1997) 3H.D. Espinosa, N.S. Brar, G. Yuan, Y. Xu, and V. Arrieta, “Enhanced
ballistic performance of confined multi-layered ceramic targets against long rod
penetrators through interface defeat,” International Journal of Solids and
Structures, 37 [36] 4893-4913 (2000). 4H.D. Espinosa, S. Dwivedi, P.D. Zavattieri, and G.Yuan, “A numerical
investigation of penetration in multilayered material/structure systems,”
International Journal of Solids and Structures, 35 [22] 2975-3001 (1998). 5Y. Li, K.T. Ramesh, and E.S.C. Chin, “Dynamic characterization of layered
and graded structures under impulsive loading,” International Journal of Solids
and Structures, 38 [34-35] 6045-61 (2001).6A.G. Evans, J. W. Hutchinson, N. A. Fleck, M. F. Ashby and H. N. G.
Wadley, “The topological design of multifunctional cellular metals,” Progress in
Materials Science, 47 [3-4] 309-27 (2001) 7H. Prielipp, M.. Knechtel, N. Claussen, S.K. Streiffer, H. Müllejans, M.
Rühle, and J. Rödel, “Strength and fracture toughness of aluminum/alumina
composites with interpenetrating networks,” Materials Science and Engineering
A, 197 [1] 19-30 (1995). 8O. Raddatz, G.A. Schneider, W. Mackens, H.Voß, and N. Claussen,
“Bridging stresses and R-curves in ceramic/metal composites,” Journal of the
European Ceramic Society, 20 [13] 2261-73 (2000). 9E.S.C. Chin, Private Communication
10M.L. Griffith and J.W. Halloran, “Freeform Fabrication of Ceramics via
Stereolithography,” Journal of the American Ceramic Society, 79 [10] 2601-608
(1996)11
E. Sachs, M.J. Cima and J.Cornie, “Three-Dimensional Printing: Rapid
Tooling and Prototypes Directly from CAD Representation”; pp. 27-47 in Solid
Freeform Fabrication Proceedings, Vol. 1. Edited by J.J. Beamen, H.L. Marcus,
D.L. Bourell, R.H. Crawford, and J.W. Barlow. University of Texas, Austin, TX,
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D.L. Bourell, H.L. Marcus, J.W. Barlow and J.J. Beamen, “Selective Laser
Sintering of Metals and Ceramics,” International Journal of Powder Metallurgy
Technology, 28 [4] 369-80 (1992) 13
J. Cesarano, “Review of Robocasting Technology,” in Proceedings of the
1998 MRS Fall Meeting, Symposium V, Solid Freeform and Additive
Fabrication, edited by D. Dimos, S.C. Danforth, and M.J. Cima, Boston, MA, pp.
133-39 (1998) 14
S. Ranngarajan, J. McIntosh, A. Bandyopadhyay, R.C. McCuiston, N.
Langrana, A. Safari, S. C. Danforth, M. Bertoldi, S. Guceri, R. B. Clancy, V.
480 Ceramic Armor Materials by Design
Jamalabad, C. Gasdaska and P. J. Whalen, “Functional Si3N4 Ceramics by Fused
Deposition: Microstructure and Mechanical Properties,” To be Submitted to
Journal of Materials Research. 15
R.C. McCuiston, B.L. Harper, S. Rangarajan, W.E. Mayo, S.C. Danforth
and C. Gasdaska, “Generation of Texture in Si3N4 made by Fused Deposition of
Ceramics (FDC) through use of -Silicon Nitride Seeds” to be submitted Journal of the American Ceramic Society
16C. Dai, G. Qi, S. Rangarajan, S. Wu, N. Langrana, A. Safari, and S. C.
Danforth, “High Quality, Fully Dense Ceramic Components Manufactured Using
Fused Deposition of Ceramics,” pp. 411-20 in Proceedings of the 7th
Solid
Freeform Fabrication Symposium, edited by D. L. Bourell, J. J. Beaman, R.H.
Crawford, H. L. Marcus and J. W. Barlow. University of Texas, Austin, TX, 1997 17
S. Rangarajan, G. Qi, N. Venkataraman, A. Safari, and S.C. Danforth,
“Powder processing, rheology, and mechanical properties of feedstock for fused
deposition of Si3N4,” Journal of the American Ceramic Society, 83 [7] 1663-
1669 (2000) 18
N. Venkataraman, S. Rangarajan, M. J. Matthewson, B. Harper, A. Safari, S.
C. Danforth, G. Wu, N. Langrana, S. Guceri, and A.Yardimci, “Feedstock
Material Property – Process Relationships in Fused Deposition of Ceramics
(FDC),” Rapid Prototyping Journal, 6 [4] 244-52 (2000) 19
E.J. Gonzalez and K.P. Trumble, “Spontaneous infiltration of alumina by
copper-oxygen alloys,” Journal of the American Ceramic Society, 79 [1] 114-20
(1996)20
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Ceramic Armor Materials by Design 481
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Ultra-Lightweight and Novel Concepts
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DEVELOPING AN ULTRA-LIGHTWEIGHT ARMOR CONCEPT
Charles E. Anderson, Jr.
Southwest Research Institute
P.O. Drawer 28510
San Antonio, TX 78228-0510
ABSTRACT
Significant reductions in armor weight have been realized over the past 30
years by the introduction of non-metallic materials (e.g., ceramics, composites,
fabrics) into armor designs. Further reduction in state-of-the-art lightweight
armors, so as to have an ultra-lightweight armor system, is a daunting challenge,
and most probably can be accomplished only by the use of materials and
geometries in novel arrangements. The process of identifying possible defeat
mechanisms and then how to exploit these mechanisms, including the
development of materials with enhanced properties, is explored.
INTRODUCTION
Armor is a synergy of mechanics and materials. I will use the term “defeat
mechanism” to denote the mechanics that the armor designer invokes to achieve a
desired affect on the threat, which for light armor is typically a small arms (rifle-
fired) bullet. A threat is characterized by its velocity, mass (inertia), geometry
(length, diameter, nose shape), and strength (flow stress and some measure of
failure, such as stain to failure). For the purposes of this paper, where we are
considering light armor, the threat is defined as the 7.62-mm armor-piercing
(APM2) bullet, shown in Fig. 1; and the 0.30-cal monolithic steel bullet
developed by Wilkins [1]. Wilkins developed the 0.30-cal bullet as a surrogate
projectile for the APM2 bullet, largely to decrease the scatter in experimental data
that resulted from fracturing of the hard steel core in the APM2 bullet. Muzzle
velocity for the bullets is 820-850 m/s. The physical characteristics of these two
bullets are summarized in Table I. The surrogate bullet has a penetration
performance that is similar to that for the APM2 bullet into hard targets.
Defeat mechanisms that might be used against an armor-piercing (AP) bullet
are shown in Table II. These defeat mechanisms are not all inclusive, and they
are often used in combination with each other. For example, tipping/rotating the
Ceramic Armor Materials by Design 485
To the extent authorized under the laws of the United States of America, all copyright interests in this publication are the propertyof The American Ceramic Society. Any duplication, reproduction, or republication of this publication or any part thereof, withoutthe express written consent of The American Ceramic Society or fee paid to the Copyright Clearance Center, is prohibited.
bullet is usually used with spaced elements with the objective of spreading the
load of the bullet onto a subsequent element.
.7840cm
3.53cm
2.74cm
.6172cm
Jacket
Point Fi l lerBase Fi l ler
Core
Fig. 1. Schematic of 7.62-mm APM2 Bullet
Table I. Physical Properties of the 7.62-mm AP Bullets
7.62-mm APM2 Bullet 7.62-mm Surrogate AP Bullet
Mass: 10.74 g Mass: 8.32 g
Length: 3.53 cm Length: 2.81 cm
Core Mass: 5.25 g Nose: 55 cone
Core Length: 2.74 cm Hardness: Rc 55
Core Hardness: Rc 62-65
Table II. Defeat Mechanisms
Deceleration (retarding force) Erosion
Obliquity Stripping the jacket
Tipping or rotating Spreading the load
Projectile fracture Blunting the nose
Spaced elements Structural response (holding the
load through a distance)
As stated in the first paragraph, armor is a synergy of mechanics and
materials. Materials are used to amplify the performance of the mechanics. And
since weight is always an issue with armor, we demand the “ultimate”
performance out of materials. The materials are pushed to their limit, that is,
failure. As Wilkins states: “The application of materials to light armor is unusual
because material properties are utilized in the region of material failure, i.e., if the
armor doesn’t fail for a given ballistic threat, it could be made lighter” [2].
These observations set up an alternative title for the paper: Why is it so
difficult to decrease the weight of a lightweight armor system? In the remainder
of the article, I will show how invoking different defeat mechanisms (often
through a change of materials) can lead to weight reductions of an armor, and also
show the difficulties inherent in achieving significant weight reductions through
evolutionary improvements in material properties.
486 Ceramic Armor Materials by Design
TRADITIONAL ARMOR
The conventional role of armor is to decelerate the projectile until it stops, i.e.,
it is defeated. The depths of penetration (DOP) as a function of impact velocity
for the AP bullet into 6061-T6 aluminum and armor steel are shown in Fig. 2.
The filled circles denote experimental data for the APM2 bullet into an aluminum
target. The lines are predictions using the Walker-Anderson penetration model
[3]. For metallic targets, semi-infinite penetration data can be used to estimate the
thickness of armor required to stop the bullet, at a specified impact velocity, by
adding approximately one bullet diameter to the semi-infinite DOP.
The bullet penetrates considerably less into armor steel than into 6061-T6
aluminum. However, the armor designer is concerned about the weight of an
armor system. The figure of merit, instead of depth of penetration (or thickness of
the target), is areal density, which is the product of the armor thickness and the
material density. The areal densities of 6061-T6 aluminum and armor steel
required to defeat the AP bullet are shown in Fig. 3. Although the bullet
penetrates considerably less into steel than into aluminum, the decrease in
penetration is not sufficient to compensate for the differences in density.
Velocity (m/s)
0 200 400 600 800 1000
Pen
etr
ati
on
De
pth
(m
m)
0
10
20
30
40
50
60
Alu
min
um
Armor S
teel
Armor Steel (eroding)
Fig. 2. DOP vs. velocity for several
metallic targets.
Are
al
De
ns
ity (
g/c
m2)
0.0
5.0
10.0
15.0
20.0
Al-
60
61-T
6
Arm
or
Ste
el
Ero
din
g S
teel
B4C
/Al
Fig. 3. Areal density required to defeat
the AP bullet at ~820 m/s.
Steels come in different strengths, and if a steel of a different strength is
substituted for the armor steel, then the penetration is changed. For example, if
mild steel is used, the areal density to stop the AP threat is approximately
22 g/cm2; if a high-hard steel is used, the areal density to stop the threat is
approximately 10 g/cm2. In general, stronger materials provide higher
decelerating forces to the penetrator. However, since there is usually a trade off
Ceramic Armor Materials by Design 487
of strength versus ductility, there is generally a limit to the strength that can be
realized. Effectively, the advantages of the increased strength are not realized
through the entire thickness of the armor element because the material fails,
generally through some damage localization process (typically involving shearing
out of an intact plug). This is the reason dual-hardness armor steel is fabricated;
the front side is made very hard, but the backside of the armor element is less
strong, but considerably more ductile.
AP bullets are very hard, and they penetrate into metallic targets in the rigid-
body penetration mode; that is, the bullets do not deform during penetration. If
the target material could be made stronger so that the bullet deforms, penetration
will decrease. If the hard steel penetrator can be made to erode—the turning of
projectile material so that there is radial flow (mushrooming), to such an extent
that the induced strains exceed the ability of the material to remain cohesive,
thereby resulting in particulation of projectile material and, as a consequence,
mass loss—then the depth of penetration is considerably reduced, as denoted by
the short dashed line in Fig. 2. Eroding penetration results in a significant
reduction in areal density, as shown in Fig. 3. To achieve erosion, a material is
required that is “harder” than the penetrator material (and so is harder than armor
steel), but is lighter than steel (so that the areal density is favorable). Ceramics
are such a material; they have very large compressive strengths, and have
densities less than that of steel. Wilkins determined that a 7.62-mm boron carbide
(B4C) ceramic tile glued to 6.35-mm 6061-T6 aluminum substrate could defeat
the AP surrogate bullet at an impact velocity of 820 m/s [4]. The areal density of
this armor is 3.62 g/cm2, which is also plotted in Fig. 3.
The response of an AP bullet against a B4C ceramic tile glued to an aluminum
(6061-T6) substrate is shown in Fig. 4. The front view shows the damage to the
ceramic, and the side view shows the deformation of the aluminum substrate
plate. Horizontal lines were drawn on the back of the substrate plate to assist in
visualizing the deformation. As can be seen, the substrate plate absorbs some of
the kinetic energy through deformation. An estimate of the kinetic energy that is
absorbed by the plate can be obtained by examining the results of VS-VR
experiments against a bare aluminum plate, where VS is the striking (impact)
velocity of the bullet, and VR is the residual velocity of the bullet after plate
perforation. The results of a number of experiments with the APM2 bullet are
shown in Fig. 5. These same data are plotted as a function of the impacting
kinetic energy (instead of VS) in Fig. 6. It is seen that the substrate material can
absorbed approximately 0.5kJ of kinetic energy.
Thus, there are requirements for armor elements with different material
properties. A hard element is needed to erode and decelerate the bullet. A ductile
element is required to capture the remnants of the eroded bullet. Materials with
different properties need to be assembled in the most advantageous way.
488 Ceramic Armor Materials by Design
(a) (b)
Fig. 4. Post-test photograph of impact of AP bullet against ceramic/aluminum
target: (a) front view of ceramic element; (b) side view of target showing
deformation of aluminum element.
VS (m/s)
0 200 400 600 800
VR (
m/s
)
0
200
400
600
800
Experiment: Normal
Experiment: Reverse
Computation: Normal
Computation: Reverse
Fig. 5. VR vs. VS for 6.35-mm-thick
6061-T6 aluminum plate.
Kinetic Energy (kJ)
0.0 0.5 1.0 1.5 2.0 2.5 3.0 3.5 4.0
VR (
m/s
)
0
200
400
600
800
Experiment: Normal
Experiment: Reverse
Computation: Normal
Computation: Reverse
Fig. 6. Residual velocity vs. striking
kinetic energy (from Fig. 5).
FURTHER WEIGHT REDUCTION
Now we want to decrease the weight of the armor further. Four possible ways
include: 1) make the front ceramic element thinner; 2) make the substrate thinner;
Ceramic Armor Materials by Design 489
3) change the substrate material; and 4) improved material properties. Each of
these will be discussed, with an emphasis on the last item.
Make the Front Ceramic Element
Thinner: Figure 7 shows the results of
decreasing the thickness of the ceramic
element. The threat defeats the target
easily, and with relatively high residual
velocity. So unless something can be
done to enhance the properties of the
ceramic (which is the fourth item),
decreasing the ceramic thickness is not
a viable option for decreasing the
weight of the system.
Make the Substrate Thinner: The
substrate material must absorb the ki-
netic energy of the residual bullet after
being decelerated and eroded by the
ceramic element. Wilkins determined
the ballistic limit, VBL, for an AD85/Al
VS
(m/s)
500 600 700 800 900 1000
VR
(m/s
)
0
100
200
300
400
500
600
700
800
5.08 mm
6.35 mm
7.62 mm
Fig. 7. Experimental VR vs. VS for AP
bullet.
substrate system as a function of ceramic thickness ( ) and substrate thickness ( )
[1]; the results for this experimental parametric study are shown in Fig. 8.
Wilkins found a significant decrease in ballistic performance for a
ceramic/aluminum substrate system when the aluminum thickness dropped below
~6 mm. He determined that this result is a consequence of the failure mode for
the substrate changing from shear plugging to petalling at 6 mm. Therefore,
the substrate cannot be made much, if any, thinner.
Change the Substrate Material: Current, state-of-the-art, lightweight armors
use a composite material in place of the aluminum substrate. Such composites
consist of Kevlar™ and polyethylene composites. Two such materials, for
example, are Gold Shield™ and Spectra Shield™, which consist of Kevlar™
fibers and polyethylene fibers, respectively, embedded in a polyethylene matrix.
The thicknesses of these composite substrate materials are considerably greater
than that of the aluminum for comparable ballistic performance. However,
because the density of the composites is considerably less than that of aluminum
(~0.90-1.2 g/cm3 vs. 2.7 g/cm
3), the overall areal density of the armor system is
decreased. In effect, this is the reverse of the aluminum-steel trade-off of density
versus strength described earlier. The areal density can be decreased by
approximately 15% using composite substrates instead of aluminum.
Improved Material Properties: Improvements in material properties can lead
to increased ballistic performance. It is not unusual to have material scientists
claim that the improvement in a material property will “naturally” result in better
490 Ceramic Armor Materials by Design
ballistic performance of the material. Since there are significant costs associated
with developing a material with enhanced properties, it is desirable to have an
estimate of the gains in ballistic performance that might be realized from such an
improvement. This is the advantage of having models, which then can be used to
make such projections. The remainder of the article will focus on the use of
models to enhance our understanding of experimental observations, and to
quantify the improvement in ballistic performance with an enhanced material
property.
Thickness, (mm)
0.0 2.5 5.0 7.5 10.0 12.5
Ba
llis
tic
Lim
it,
VB
L (
m/s
)
0
200
400
600
800
1000 = 8.64 mm
= 7.87 mm
= 6.35 mm
= 4.06 mm
Fig. 8. Ballistic limit velocity as a function of ceramic and substrate
thickness for AD85 (Al2O3)/6061-T6 Al (from Wilkins [1]).
ANALYTICAL AND COMPUTATIONAL MODELING
We would like to use the results of modeling to guide armor development. In
particular, we would like to investigate, and quantify, the advantage of improved
material properties. In order for modeling to be useful for this endeavor, it must
be demonstrated that the modeling captures the essential features of observed
phenomena, and that the modeling provides reasonable agreement with
experimental data. It is not necessary for the model to reproduce exactly
experimental results, but it is necessary that the model be sufficiently accurate so
that it can predict the correct trends. This is why the first requirement is
necessary—that the model captures the essential features of observed
phenomena—because model parameters can be tuned to provide good agreement
with experimental results, but not have the correct mechanics and physics.
Figure 9 shows flash radiographs of the APM2 bullet, 15.3 s and 20.7 s
after impact, against a 7.62-mm-thick B4C tile backed by a nominal 6.35-mm-
Ceramic Armor Materials by Design 491
thick, aluminum (6061-T6) substrate plate. The impact velocity was approxi-
mately 820 m/s for these experiments. The bullet is “dwelling” at the surface of
the ceramic (not penetrating) in the first image; by approximately 20 s, the
integral strength of the ceramic no longer can support dwell, and the bullet is
penetrating (the right image).
(a) 15.3 s (b) 20.7 s
Fig. 9. Flash radiographs of the APM2 bullet impacting a B4C/Al target.
A simple analytical model of dwell has
been developed; the idealized model is
shown in Fig. 10. The governing equations
are shown below the figure, where p is the
projectile density, v is the tail velocity, is
the current length of the projectile, Y
λp is the
projectile flow stress, u is the penetration
velocity, and t denotes time. The first equa-
tion describes deceleration of the tail, and the
second equation describes the shortening
(and thus mass loss) of the projectile. The
third equation is the statement of the assump-
tion that the penetration velocity, u, is zero.
These equations can be solved explicitly.
The solutions for the surrogate AP projectile,
at an impact velocity of 820 m/s, are shown
������������������������������������������������������������������������������������������������������������������������������������������������������������������������������������������������������������������������������������������
VV
0
)v(
vpp
u
udt
d
Ydt
d
λ
λ
Fig. 10. Analytical model for
dwell.
in Figs. 11-13. The results of deceleration of the bullet as a function of time are
shown in Fig. 11. The length of the bullet decreases because of erosion, with an
attendant loss of mass, Fig. 12. Mass loss, initially, is quite small because of the
pointed noise.
The kinetic energy of the bullet as a function of time is plotted in Fig. 13. The
percentage of kinetic energy lost to erosion, and that lost to deceleration, can be
492 Ceramic Armor Materials by Design
separated. It is seen that each of the “defeat” mechanisms contributes to a
significant loss in kinetic energy of the AP bullet. Although the model provides
an idealized description of dwell, it permits a quantification of the advantages if
dwell can be extended for a few additional microseconds. For example, the flash
radiograph in Fig. 9(a) was taken at 15.3 s after impact. At this time, the bullet
has lost 44% of its initial kinetic energy. If dwell could have been made to extend
to 20.7 s, Fig. 9(b), then the kinetic energy would have decayed to
approximately 23%, a significant decrease in kinetic energy.
Time ( s)
0 5 10 15 20 25 30
% K
ineti
c E
nerg
y
0
10
20
30
40
50
60
70
80
90
100K
ineti
c E
nerg
y (
kJ)
0.0
0.5
1.0
1.5
2.0
2.5
KE with Mass Loss
(Erosion Only)
Erosion & Deceleration
KE lost due
to erosion
KE lost due
to velocity decay
Time ( s)
0 5 10 15 20 25 30
Ve
loc
ity (
m/s
)
0
200
400
600
800
1000
Fig. 11. Velocity vs. time for
dwelling AP bullet.Fig. 12. Length and mass vs. time for
dwelling AP bullet.
Time ( s)
0 5 10 15 20 25 30
Mass (
g)
0.0
1.0
2.0
3.0
4.0
5.0
6.0
7.0
8.0
9.0
Len
gth
(m
m)
0.0
5.0
10.0
15.0
20.0
25.0
30.0
Fig. 13. Kinetic energy vs. time for dwelling AP bullet.
Ceramic Armor Materials by Design 493
The analytical dwell model is useful in quantifying the effects of dwell, but it
cannot predict if dwell will occur, and if it does, for how long it will last. To
make this prediction, we need to turn to a computational model. Anderson and
Walker [5] modified a computational ceramics model developed by Wilkins [1-2].
The modified model, implemented into the wavecode CTH [6], has 5 model
constants: intact (compressive) strength of the ceramic, tensile strength of the
ceramic, the slope and cap of a Drucker-Prager yield surface for the damaged
(comminuted) ceramic, and a constant that governs the speed of damage from
intact to comminuted ceramic (a fraction of the shear wave velocity). A damage
parameter, f, is defined: 0f implies intact ceramic; 1f denotes
completely failed ceramic within a computational cell. Failure of a computational
cell is initiated when the calculated tensile stress exceeds the material tensile
strength, subject to the condition that a cell is next to a cell that has failed, 1f ,
or is next to a material interface or free surface. Once damage is initiated, the
strength of the computational cell goes from that of intact material to that of the
comminuted material at the prescribed damage rate (the fifth parameter). All
parameters but the last are determined from independent laboratory experiments;
the last parameter was calibrated to achieve the correct residual length (LR) of
recovered cores from the APM2 bullet. (Thus, the intact ceramic is modeled
elastic-plastic until failure; thereafter, the failed or comminuted ceramic strength
follows a Drucker-Prager constitutive relationship. The metallic elements—
projectile and substrate materials—are modeled as elastic-plastic, with strain
hardening, strain rate and temperature effects. All materials are considered
isotropic.) The modified model reproduces quite accurately a wide variety of
experimental results of impact into thin ceramic tiles, including the phenomenon
of dwell.
The nose and tail velocities of one such simulation, of the AP surrogate bullet
into 7.62-mm B4C/6.35-mm 6061-T6 Al, is shown in Fig. 14. Wilkins showed
that this armor configuration stopped the bullet; the simulation is in agreement
with experiment. The simulation results indicate that dwell lasts for
approximately 20 s, and that the bullet penetrates as a rigid body (nose and tail
velocities the same) after 26 s. It is interesting to note that the kinetic energy
remaining after 23 s of dwell (Fig. 13), is approximately 0.5 kJ, the same energy
that can be absorbed by 6.35 mm of aluminum (Fig. 6).
The analytical dwell model and the computational model provide an
explanation for the “sudden” drop off in performance of our ceramic armor
system as the ceramic tile is made thinner, Fig. 7. If it is assumed that dwell does
not last as long if the ceramic tile is made thinner, then there is less deceleration
of the bullet, less erosion of the bullet, and the resulting kinetic energy of the
“remnant” bullet when it reaches the substrate is considerably higher than 0.5 kJ
(see Figs. 11-13). Evidence of less erosion as the ceramic tile is made thinner is
494 Ceramic Armor Materials by Design
seen in recovered cores, Fig. 15. However, it turns out to be even worse than
simply a decrease in dwell time. According to the computational model, dwell
hardly persists for the thinner tiles, Fig. 16. There is a pseudo-dwell period where
the penetration velocity is relatively small, but an almost zero penetration velocity
is not predicted. The reason for this will be discussed a little later.
Time ( s)
0 10 20 30 40 50 60 70 80 90 100
Vel
oci
ty (
m/s
)
0
100
200
300
400
500
600
700
800
900
Tail
Nose
Fig. 14. Nose and tail velocities vs. time from numerical simulation of the
surrogate bullet against 7.62-mm B4C/6.35-mm 6061-T6 Al.
VS (m/s)
600 700 800 900 1000
LR
(m
m)
0
5
10
15
20
25
5.08 mm
6.35 mm
7.62 mm
Experiment--5.1 mm
Experiment--6.4 mm
Calculation--5.08 mm
Calculation--6.35 mm
Experiment--7.62 mm
Fig. 15. Residual length of AP cores.
Time ( s)
0 10 20 30 40 50 60 70
Vel
oci
ty (
m/s
)
0
100
200
300
400
500
600
700
800
900
5.08 mm
6.35 mm
7.62 mm
Fig. 16. Nose and tail velocities for
different ceramic tile thickness.
Ceramic Armor Materials by Design 495
IMPROVED MATERIAL PROPERTIES
The models will now be used to quantify the gains that might be expected if a
boron carbide ceramic could be made with improved material properties. Within
the context of the computational model, the two parameters that might be
improved through changes in processing are the compressive and tensile
strengths. Simulations indicate that increasing the compressive strength of the
ceramic does not substantially change the ballistic performance of the ceramic.
This implies that the ceramic is already sufficiently hard to erode the AP bullet. It
might be expected, however, that since the model is a tensile-failure-initiation
model, that improving the tensile strength of the ceramic will improve ballistic
performance.
Computational results, using the
measured tensile strength of B4C ( f =
0.3 GPa), are compared to experi-
mental results in Fig. 17 for a tile
thickness of 5.08 mm (over a 6.35-mm
Al substrate). Good agreement is seen
for VR vs. VS. So additional simula-
tions were performed where the tensile
strength was increased by a factor of 3
and 5, to 0.9 and 1.5 GPa, respectively.
The results are shown for a 5.08-mm-
thick tile in Fig. 18. Even with a five-
fold increase in the tensile strength, the
residual velocity for an 850-m/s-impact
velocity is over 300 m/s. A strength of
1.5 GPa is equivalent to the flow stress
of a hard armor steel, so it is not clear
VS (m/s)
500 600 700 800 900
VR (
m/s
)
0
100
200
300
400
500
600
700
f = 0.3 GPa
Fig. 17. Computational and
experimental VR vs. VS (5.08-mm B4C)
that such a ceramic can even be fabricated. Even if it such a “new” ceramic could
be fabricated, the overall areal density would change from 3.62 g/cm2 (7.62-mm
B4C) to something slightly greater than 3.0 g/cm2 (5.08-mm B4C), a change of
only 18% for a fivefold increase in the tensile strength!
The simulation results were analyzed to determine why such a dramatic
increase in a physical property has so little influence on ballistic performance.
The minimum principle stress throughout the ceramic tile was plotted at a number
of times after impact. These plots show that tensile stresses within the ceramic
element exceed 1.0 GPa in the entire volume under the penetrator during the first
few microseconds, with some areas having tensile stresses in excess of 1.5 GPa
(for the 5.08-mm-thick tile). Thus, the problem is that the impact event is so
severe that the material is simply “overwhelmed” by the dynamics. As the
ceramic tile is made thinner, the ability to resist tensile stresses decreases
496 Ceramic Armor Materials by Design
nonlinearly (bending stiffness is proportional to the thickness cubed). Therefore,
increased tensile strengths of 3 to 5 times the current material property value are
not sufficient to compensate for the increased tensile stresses generated from
impact.
VS (m/s)
500 600 700 800 900
VR (
m/s
)
0
100
200
300
400
500
600
700
f = 0.3 GPa
f = 0.9 GPa
f = 1.5 GPa
Fig. 18. Effects of increased tensile
strength: 5.08-mm B4C
VS (m/s)
500 600 700 800 900
VR (
m/s
)
0
100
200
300
400
500
600
700
f = 0.3 GPa
f = 0.9 GPa
5.08 mm
6.35 mm
Fig. 19. Effect of increased tensile
strength: 5.08-mm & 6.35-mm B4C
Nevertheless, failure time through the ceramic element is increased by the
increased f, as can be inferred from the decrease in residual velocity. That is, the
armor system can be made lighter using the “improved” material. Predictions of
VR vs. VS for a 6.35-mm-thick tile are shown in Fig. 19. VR, for an impact
velocity of 850 m/s, is approximately 200 m/s for f = 0.9 GPa. Experience has
shown that when VR’s are ~200 m/s or lower—because of the steepness of the
VR-VS curve near the ballistic limit—the armor system is approximately at the V50
design. So the increased fracture strength does have an effect on ballistic
performance, but the effect is not nearly as large as one might have thought based
on the significant increase in f. The increase of f from 0.3 GPa to 0.9 GPa
results in a decrease in the weight of the armor system from 3.62 g/cm2 to
approximately 3.30 g/cm2. Unfortunately, it has taken a significant improvement
in a material property to realize a 9% decrease in areal density.
SUMMARY AND CONCLUSIONS
Light armor is a synergy of mechanics and materials. Because the armor
designer is demanding the ultimate performance out of the materials that are being
used, the performance of lightweight armor is “precipitous,” i.e., a very small
change in geometry (for example, a small decrease in the thickness of an armor
Ceramic Armor Materials by Design 497
element) and the armor is defeated quite easily. This “precipitous” behavior
makes it difficult to decrease the weight unless the operative mechanics (defeat
mechanisms) are changed (e.g., adding erosion to deceleration), or unless a
material is changed (e.g., changing the substrate from aluminum to a composite).
Thus, as shown by our example of an improved ceramic (an increase in the
tensile strength of the ceramic), evolutionary changes in material properties result
in incremental changes in ballistic performance, and incremental decrease in
weight. A significant increase in ballistic performance (i.e., a significant decrease
in weight) requires an advance defeat mechanism (or invoking several defeat
mechanisms), and/or a revolutionary advance in materials.
ACKNOWLEDGEMENT
The author would like to thank Dr. Steve Wax of DARPA and Mrs. Janet
Ward of the U. S. Army Soldier Systems Command for their support and
suggestions in the preparation of this paper. The author would also like to thank
Mr. Dick Sharron (SwRI) for his assistance in running of the numerical
simulations. This work was funded under contract DAAD16-00-C-9260.
REFERENCES 1M. L. Wilkins, “Mechanics of Penetration and Perforation,” Int. J. Engng.
Sci., 16(11), 793-807, 1978. 2M. L. Wilkins, “Third Progress Report of Light Armor Program,” UCRL-
50460, Lawrence Livermore Laboratory, Livermore, CA, July 1968. 3J. D. Walker and C. E. Anderson, Jr., “A Time-Dependent Model for Long-
Rod Penetration,” Int. J. Impact Engng., 16(1), 19-48, 1995. 4M. L. Wilkins, R. L. Landingham, and C. A. Honodel, “Fifth Progress Report
of Light Armor Program,” UCRL-50980, Lawrence Livermore Laboratory,
Livermore, CA, 1970. 5C. E. Anderson, Jr. and J. D. Walker, “Ceramic Dwell and Defeat of the 0.30-
Cal AP Projectile,” 15th
U.S. Army Symp. on Solid Mech., Myrtle Beach, SC,
April 12-14, 1999. 6J. M. McGlaun, S. L. Thompson, and M. G. Elrick, “CTH: A Three-
Dimensional Shock Wave Physics Code,” Int. J. Impact Engng., 10, 351-360,
1990.
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To the extent authorized under the laws of the United States of America, all copyright interests in this publication are the propertyof The American Ceramic Society. Any duplication, reproduction, or republication of this publication or any part thereof, withoutthe express written consent of The American Ceramic Society or fee paid to the Copyright Clearance Center, is prohibited.
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NOVEL IDEAS IN MULTI-FUNCTIONAL CERAMIC ARMOR DESIGN
Sia Nemat –Nasser*, Sai Sarva, Jon B Isaacs and David W Lischer
Center of Excellence for Advanced Materials
Department of Mechanical and Aerospace Engineering
University of California, San Diego
La Jolla, CA 92093-0416
ABSTRACT
Ceramics such as Al2O3, SiC, TiB2 and B4C have been used in integrated
armor for over a decade and are an excellent prospect for the next generation
multi-functional armor systems. It is necessary to incorporate novel ideas in
ceramic armor design so as to develop improved armor with minimal added mass.
Preliminary research has demonstrated that the defeat capability of ceramic armor
tiles could be considerably improved by tightly wrapping them in a thin
membrane of suitable tensile strength. In the present paper we present some
recent experimental results relating to the effect of thin membranes attached to the
front face of Al2O3 armor tiles, on their ballistic performance. The experiments
were conducted to study the comparative effect of several front-face materials,
such as glass-fiber tape, E-glass/epoxy pre-preg, Carbon-fiber/epoxy pre-preg and
Ti-3%Al-2.5%V alloy. Tungsten heavy alloy was used as the projectile material.
It was observed that confinement by a thin layer of E-glass/epoxy pre-preg
resulted in a nearly 20% improvement in the ballistic efficiency for a mere 2.5%
increase in areal density. The improvement in ballistic efficiency is accompanied
by an altering of the failure mechanisms. High-speed photography and flash
radiography techniques have been used to gain insight into the mechanisms that
may be responsible for this improvement.
INTRODUCTION
The next generation armor systems require integration of several attributes
within hybrid structures, which can be accomplished through introduction of
novel concepts in the materials-structural design. These attributes may include
great agility, effective communication, and controlled signature. New
Corresponding author: [email protected] (858) 534-4914, Fax: (858) 534 2727
Ceramic Armor Materials by Design 511
To the extent authorized under the laws of the United States of America, all copyright interests in this publication are the propertyof The American Ceramic Society. Any duplication, reproduction, or republication of this publication or any part thereof, withoutthe express written consent of The American Ceramic Society or fee paid to the Copyright Clearance Center, is prohibited.
materials/structures must therefore be created in such a manner that they are light-
weight, impact resistive, have structural integrity and at the same time can have
signature management and controlled communication capabilities. This can be
achieved by incorporating periodic arrays of thin conductor wires1, exhibiting the
desired electro-magnetic response into high strength, low-density hybrid
composites. Extensive research has been conducted to increase the ballistic
efficiency to areal density ratio of ceramics through various techniques. A number
of researchers have studied the effect of confinement of ceramics on their ballistic
performance and failure modes. Shockey et al.2 studied the failure
phenomenology of confined ceramics under rod impact. They concluded that the
key processes are crushing of the ceramic and the subsequent flow of fine
fragments lateral to and opposite to the direction of impact. Woodward et al.3
studied the perforation of confined and unconfined ceramic targets by pointed and
blunt projectiles. It was observed that front confinement of ceramic results in a
greater overall fragmentation. However, their experiments suggest that less
amount of very fine ceramic powder may form in the confined case as compared
to the unconfined target. Anderson and Morris4 have studied the effect of
projectile diameter on its erosion for thick (~ 4 cm) Al2O3 tiles under lateral and
rear confinement. They also observed that for constant-mass projectiles, longer
rods erode more than shorter rods for the same ceramic thickness.
When projectiles impact ceramic targets, a pulverized zone is formed ahead of
the projectile head due to intense stress conditions. Understanding of the failure
mechanisms resulting in this pulverization is important for developing improved
models and for designing better armor systems. Curran et al.5 present a
micromechanical model for comminution and granular flow of ceramics under
impact. Cortes et al.6 have numerically modeled the impact of ceramic-composite
armor. They present a constitutive model for finely pulverized ceramic taking into
account internal friction and volumetric expansion. Grace and Rupert7 have
incorporated models of Curran et al.5 and Cortes et al.
6 to analyze long rods
penetrating ceramic targets at high velocities. McGinn7
et al.8 have
microscopically studied the deformation and comminution of shock loaded Al2O3
to understand the failure mechanisms that produce this pulverized zone, often
referred to as the ‘Mescall zone’.
Recently, McGee et al.9 studied the effect of thin membrane wrapping on the
defeat capability of Al2O3 and SiC ceramic armor tiles. It was observed that
tightly hand-wrapping the tiles in commercially available Scotch glass fiber tape
improves the ballistic efficiency by nearly 20%. It was also observed that this
improvement is mainly a result of impact-face constraint that the tape provides,
and that the back-face constraint had little (if any) effect on the ballistic
efficiency.
512 Ceramic Armor Materials by Design
Following these observations, further experiments were conducted to
investigate the effect of the front-face-attached membrane on the failure
mechanisms and the projectile-target interaction. Also, the material and the
thickness of the front-face-attached membrane were varied to observe the
resulting effects on the ballistic efficiency. The experimental results are discussed
in the present paper together with some numerical simulations, leading to some
tentative conclusions on the potential factors that may be involved in this process.
EXPERIMENTAL PROCEDURES
Gas-Gun
A single stage gas-gun is used to launch the projectile. Helium is the
driving gas. The barrel diameter is 2.54 cm and its length is 4.8 m. Two velocity
sensors at the muzzle end of the barrel are used to measure the intial velocity of
the projectile. The sensors also trigger the high-speed camera and flash X-ray
heads. The gas-gun can launch a 17 gm sabot-projectile assembly at up to about
1100 m/s. The gas gun is operated with two different configurations of target
assembly depending on the nature of data of interest.
The stripped-sabot configuration: An Aluminum sabot carries the projectile
through the barrel. Prior to impact, the sabot is stripped by means of a maraging
steel stripper. After penetration, the projectile erodes and its velocity reduces. The
residual velocity is measured by means of residual velocity sensors. The
projectile is recovered from paper stacks, which act as momentum dump and the
residual mass is measured. Ballistic performance is evaluated by comparing the
kinetic energy of the residual rods. See Fig. 1.
The unstripped-sabot configuration: Using a sabot-stripper creates sabot debris
during the stripping process. This debris interferes with high-speed photography.
Hence, tests were also conducted without the stripper and the residual velocity
sensors. This configuration provides immaculate imagery of the initial stages of
the impact phenomenon and helps study ejecta characteristics. However, the time
window of data acquisition is limited to until the sabot interferes with the
penetration process.
Ceramic Armor Materials by Design 513
Figure 1. Stripped-sabot configuration for ballistic tests
High-speed Photography
The Hadland Imacon 200 high-speed image acquisition system was used to
study the ultra high-speed phenomenon of ballistic penetration. The camera can
be programmed to record a sequence of separate images at prescribed time
intervals. A sixteen-channel camera was used. Images were acquired from a point
of view normal to the path of the projectile.
Flash Radiography Procedures
During ceramic penetration, fine pulverized ceramic powder is ejected from
the front and rear surfaces of the tile. This obscures the view of projectile-target
interaction and the flow of eroded particles. An experimental set-up for flash
radiography provides dynamic, real time images of the projectile penetrating the
ceramic. Two 100 kV heads were used. Two configurations were used. In the
inclined X-ray configuration, as seen in Fig. 2, the X-ray heads are placed
inclined to the path of the projectile. This reduces the ceramic cross-section that is
pierced by the X-rays. This configuration helps study the interior of the target and
hence the target-projectile interaction during penetration. In the edge-on
514 Ceramic Armor Materials by Design
configuration, the X-ray heads (see Fig. 2) are moved so that they are orthogonal
to the path of the projectile. Since, the target thickness is large this configuration
does not reveal the interior. It helps study the flow of rod erosion products
emerging from the front surface.
Inclined X-ray Edge-on X-ray
Figure 2. Flash radiography configurations
Target Material
Coors Al2O3 AD995 CAP3 armor grade tiles were used. These are 99.5%
purity tiles of 10.16 cm 10.16 cm 1.27 cm dimensions. The areal density of
the ceramic tiles is 4.98 gm/cm2.
2.5 Projectile Material
WHA (93% W, ~5% Ni, ~2% Fe) manufactured by Hogen Industries was
used. The projectiles were flat-ended cylinders of diameter 6.14 mm and length
20.86 mm. Also, WHA (93%W, ~5%Ni, ~2%Fe) procured from ARL was used
for flash radiography studies.
Membrane application techniques
Scotch 893 Glass Fiber tape: Commercially available Scotch 893 glass fiber tape
was used to hand-wrap the ceramic tiles. Scotch fiber tape has a tensile strength of
525 N/cm. It is 0.15 mm thick. Elongation is approximately 4.5%. Eight layers of
fiber tape were hand-wrapped on tiles and then the back-face tape was cut out so
that only the front-face and edges of the ceramic tiles were taped. The glass fibers
on the cellophane tape run uni-directionally. Hence, the orientation of the tape
Ceramic Armor Materials by Design 515
was alternated after every two layers (02/902/02/902). Taping the ceramic tile
increases its areal density from 4.98 to 5.31 g/cm2.
Ti-3%Al/2.5%V sheets: Ti-3/2.5 alloy sheets of 0.127 mm, 0.254 mm and 0.508
mm were used. The sheets were bonded to the front-face of the ceramic tiles
using Loctite 312 super glue. The tensile strength of Ti-3/2.5 is approximately
620 MPa. Elongation is approximately 15%. The areal densities of ceramic tiles
with 0.127mm, 0.254 mm and 0.508 mm Ti-3/2.5 sheets are 5.036 gm/cm2, 5.093
gm/cm2 and 5.207 gm/cm
2 respectively.
E-glass/Epoxy pre-preg: E-glass/Epoxy pre-preg (BT-250E-1) manufactured by
Bryte technologies Inc. was used. The E-glass reinforcement has a cross weave
and the overall tensile strength is 434 MPa. The pre-preg was pressed onto the
front surface of the ceramic tiles and cured at 250o F (121
o C), in a hot press.
Samples with one and three layers of pre-preg were prepared. The areal densities
are 5.019 gm/cm2 and 5.099 gm/cm
2 respectively.
Carbon fiber/Epoxy pre-preg: Carbon-fiber/Epoxy pre-preg (BT-250E-1), also
manufactured by Bryte technologies was used. The Carbon (Graphite)
reinforcement also has a cross weave and the overall tensile strength is 669 MPa.
Samples were prepared using techniques similar to those used for E-glass/Epoxy
prepreg. Samples with one and three layers of pre-preg have areal densities of
5.017 gm/cm2 and 5.083 gm/cm
2 respectively.
EXPERIMENTAL RESULTS
WHA projectiles were used to impact Al2O3 tiles at 900 m/s. The velocity was
well above the ballistic limit (V50) of the Al2O3 tiles and was maintained the same
for all the tests. Bare tiles and tiles with front-face fiberglass tape, Ti-3/2.5,
Carbon fiber/Epoxy pre-preg, or E-glass/ Epoxy pre-preg membrane of various
thicknesses, were studied.
Ballistic performance
Tests conducted with the stripped sabot configuration help to understand the
effect of impact-face constraint on the ballistic performance. The projectiles
weighed about 10.6 gm. The initial velocity of the projectile, measured by the
velocity sensors was used to calculate the initial kinetic energy. The residual
velocity sensors measured the exit velocity of the projectiles, after penetration.
The eroded projectiles were recovered from the paper stacks and weighed. The
residual kinetic energy was calculated. The ballistic performance was evaluated
by determining the kinetic energy fraction, defined by fKE = residual kinetic
energy/ initial kinetic energy. Some of the results are shown in Fig. 3.
516 Ceramic Armor Materials by Design
As can be seen, the fKE for bare tiles is approximately 0.35. From the test
results for the E-glass/Epoxy, Carbon-fiber/Epoxy, and Ti-3/2.5, it is observed
that fKE tends to diminish with increasing thickness of the membrane layer. The
fKE for a three layered E-glass/Epoxy sample is approximately 0.12, about the
third of that for the bare tiles. This is a nearly 23% improvement in the ballistic
efficiency for a mere 2.5% increase in the areal density. It is also observed that
glass-fiber tape improves the ballistic efficiency substantially. However, the areal
density is increased by nearly 7%, mainly as a result of the cellophane content. It
is expected that after a certain critical thickness for the front-face membrane,
there will be a gradual reduction in the resulting improvement due to the
constraint effect.
0.05
0.12
0.19
0.26
0.33
4.95 5 5.05 5.1 5.15 5.2 5.25 5.3 5.35
Areal density ( gm/cm^2 )
Kin
eti
c E
nerg
y F
racti
on
( Unconfined Alumina )
( 5 mil Ti )
(10 mil Ti)
(20 mil Ti)
(Glass Fiber Tape)
8 layers - 48 mil
E Glass/Epoxy prepreg
1 layer - 8 mil
E Glass/Epoxy prepreg
3 layers - 24 mil
Carbon fiber/Epoxy
1 layer
Carbon fiber/Epoxy
3 layers
Figure 3. The effect of front-face constraint on the ballistic performance of
allumina tiles
Table 1. lists the residual velocity and residual mass measurements of the
projectiles. As can be seen from the Ti-3/2.5, Carbon-fiber/Epoxy, and E-
glass/Epoxy tests, the residual velocity is decreased by more than 100 m/s for the
front-face-constraint samples. However, no strong correlation is yet observed
between the residual velocity and increasing thickness of the membrane. It is
observed that increasing the thickness of the membrane results in an increase in
erosion. Hence, preliminary observations suggest that increasing thickness and
hence the ensuing increase in tensile strength of the impact-face membrane,
increases the erosion of the projectile. Further tests are needed to isolate the
Ceramic Armor Materials by Design 517
effects of key material properties such as tensile strength, stiffness, and elongation
of the front-face membrane on the residual velocity and erosion of the projectiles.
Constraning
membrane material
Initial
velocity (m/s)
Initial
Mass (gms)
Residual
velocity (m/s)
Residual
mass (gms)
Unconfined 903.9 10.708 682.0 6.421
Unconfined 900.7 10.658 671.0 6.489
Glass fiber tape 897.5 10.582 545.5 5.706
Glass fiber tape 900.7 10.582 563.7 4.955
Ti – 0.127 mm 887.5 10.668 624.4 6.309
Ti – 0.127 mm 912.0 10.662 584.7 5.008
Ti – 0.254 mm 900.7 10.570 636.7 3.567
Ti – 0.254 mm 894.4 10.582 561.7 4.429
Ti – 0.508 mm 891.2 10.634 616.0 2.328
Carbon – 1 lyr 894.4 10.647 632.5 6.791
Carbon – 1 lyr 905.5 10.671 633.6 5.846
Carbon – 3 lyrs 892.3 10.660 540.8 4.805
Carbon – 3 lyrs 891.2 10.663 538.0 4.339
E-glass – 1 lyr 892.7 10.622 593.4 5.721
E-glass – 1 lyr 900.7 10.610 527.5 4.813
E-glass – 3 lyrs 907.1 10.656 517.0 3.808
E-glass – 3 lyrs 864.1 10.615 532.1 4.253
Table 1. The effect of various constraining materials on the residual velocity and
the residual mass of the projectile
High speed photography results
Front face: Fig. 4 shows the initial stages of an unstripped-sabot test for a bare
sample. The ceramic ejecta can be seen ejecting from the front surface. Soon
after impact, a pulverized zone (the Mescall zone) is formed ahead of the
projectile due to intense stress conditions. The ejection process clears the
pulverized ceramic away to accommodate the penetration of the projectile. A
significant portion of the kinetic energy of the projectile is transferred to the
ejecta. As can be seen, the ejecta for a bare sample is radially disperse and conical
in shape. Fig. 5 shows the initial stages of an unstripped-sabot test for a
constrained sample. The flow of ejecta particles is much more acute and
518 Ceramic Armor Materials by Design
cylindrical in nature. The displacement measurement tools of the Imacon 200
software were used to calculate the velocities of these ejecta. It was observed that
during the initial stages, the ejecta velocity for samples with front-face membrane
was nearly 40% higher than that of the corresponding bare samples. The higher
kinetic energy associated with ejecta signifies reduced residual kinetic energy for
the projectile.
(1 ) (5 )s s
(9 ) (13 )s s
Figure 4. Initial stages of impact of a bare tile
Ceramic Armor Materials by Design 519
(0 ) (4 )s s
(8 ) (12 )s s
Figure 5. Initial stages of impact of E-glass/Epoxy constrained tile
(10 ) (15 )s s
(20 ) (25 )s s
Figure 6. Back face displacement of a bare tile
520 Ceramic Armor Materials by Design
(10 ) (15 )s s
(20 ) (25 )s s
Figure 7. Back-face displacement of an E-glass/Epoxy constrained sample
Back face displacement: The projectile’s travel velocity and the rate of its
erosion govern its penetration rate. The back-face displacement gives a good
indication of the penetration rate. Figs. 6 and 7 show the back-face displacement
of a bare and a front-face constrained sample, respectively. It can be seen that the
back-face displacement is delayed by nearly 15 for the constrained sample.
This implies increased erosion and/or reduction in velocity.
s
Flash Radiography
Fig. 8 compares the X-ray images for bare and front-face-membrane
constrained samples. The edge-on X-rays indicate that the eroded projectile
particles for the constrained sample, exhibit a more oblique flow as compared to
that for the bare sample. The inclined X-rays provide a view of the interior during
penetration. The projectiles deform by mushrooming and shearing of its tip,
indicating ductile nature of its failure. The projectile for the constrained sample
exhibits a larger mushroom head. This confirms the increased penetration
resistance and erosion of the projectile. The more oblique flow of the eroded
projectile particles for a constrained sample, observed in the edge-on X-rays, is a
result of greater mushrooming.
Ceramic Armor Materials by Design 521
(7 ) (15 ) (8 )s s s
Edge-on X-rays Inclined X-raysBare tile
(7 ) (15 ) (9 )s s s
Edge-on X-rays Inclined X-raysConstrained tile
Figure 8. Flash radiography results
CONCLUSIONS
Al2O3 tiles when impacted by WHA projectiles fail through a complex
combination of processes resulting from the shock-wave propagation and
reflection. These processes include fragmentation and formation of radial and
circumferential macro-cracks, pulverization of the ceramic into fine powder, and
ejection of the fine powder from front and rear surfaces. The WHA projectile
undergoes deformation and erosion. It is expected that the morphology of the
pulverized ceramic fragments and its flow characteristics govern the penetration
resistance of the ceramic tiles. Hence, it is important to understand the underlying
mechanisms producing the pulverization of ceramics. Preliminary numerical
simulations10
on DYNA2D (a two-dimensional hydrodynamic finite element
code) indicate that release waves emanating from the projectile edges reduce the
pressure and increase the shear stress at a distance equal to the projectile diameter,
ahead of the projectile. The resulting stress condition is highly conducive to the
pulverization of ceramic11, 12
, See Fig. 9.
522 Ceramic Armor Materials by Design
Figure 9(a) Computational grid displaying the geometry
0 0.1 0.2 0.3 0.4 0.5 0.6 0.7 0.8 0.9 1 1.1 1.2 1.3 1.40
0.1
0.2
0.3
0.4
0.5
0.6
0.7
0.8
0.9
1
1.1
1.2
x (cm)
y (c
m)
Maximum Shear Stress (GPa)
15.8755
14.4
323
12.9
891
12.9891 11.545
8
10.1
026
8.65
94
7.216
2
5.772
9
4.3297
2.8865
1.4432
4.32
97
5.7729
Ceramic Armor Materials by Design 523
0 0.1 0.2 0.3 0.4 0.5 0.6 0.7 0.8 0.9 1 1.1 1.2 1.3 1.40
0.1
0.2
0.3
0.4
0.5
0.6
0.7
0.8
0.9
1
1.1
1.2
x (cm)
y (c
m)
Pressure (GPa)
30.39
01
22.31
14
14.232
68.
8468
6.15
393.
461 0.76811
0.76
811
3.461
6.15
39
0.76811
Figure 9 (b) and (c). Contours of maximum shear stress and constant ressure in
an Al2O3 tile 0.83 s after being impacted by a 6.35 mm WHA projectile 9
High-speed photographs indicate that the front face confinement of Al2O3 tiles
vastly alters the flow of the pulverized ceramic that is ejected out. The front-face
ejecta from a bare tile is radially disperse and conical. For a constrained tile the
ejecta flow is more acute and cylindrical. Also the ejecta velocity for a
constrained tile is nearly 40% higher. Flash radiography indicates that
constraining the ceramic tile results in a much greater mushrooming and erosion
of the projectile. The greater erosion and reduced velocity of the projectile are
also manifested in the form of a significant delay in the back-face displacement of
the ceramic tile. Experiments indicate that impact-face restraint by fiber
reinforced polymer results in a substantial improvement in the ballistic efficiency.
Thin layers of E-glass/Epoxy improve the ballistic efficiency by nearly 20% for
an increase in areal density of 2.5%. Further research is needed to study the effect
of front-face membrane of other materials, so as to isolate the key material
properties governing the improvement in ballistic efficiency.
ACKNOWLEDGEMENT
The reported work was supported by US Army Research Office under contract
No ARO DAAH04-96-1-0376, to University of California at San Diego
524 Ceramic Armor Materials by Design
REFERENCES1D.R. Smith, D.C. Vier, W. Padilla, S. C. Nemat-Nasser and S. Schultz, “Loop
wire medium for investigating plasmons at microwave frequencies,” Appld. Phys.
letters 75[10], 1425-1427, (1999)2D.A. Shockey, A.H. Marchand, S.R. Skaggs, G.E. Cort, M.W. Burkett and R.
Parker, “Failure phenomenology of confined ceramic targets and impacting rods,”
Int. J. Impact Engng. 9[3], 263-275 (1990) 3R.L. Woodward, W.A. Gooch Jr, R.G. O’Donnell, W.J. Perciball , B.J.
Baxter and S.D. Pattie, “A study of fragmentation in the ballistic impact of
ceramics,” Int. J. Impact Engng. 15[5], 605-618 (1994)4C.E. Anderson, Jr and B.L. Morris, “The ballistic performance of confined
Al2O3 ceramic tiles,” Int. J. Impact Engng. 12[2], 167-187 (1992) 5D.R. Curran, L. Seaman, T. Cooper and D.A. Shockey, “Micromechanical
model for comminution and granular flow of brittle material under high strain rate
application to penetration of ceramic targets,” Int. J. Impact Engng. 13[1], 53-83,
(1993)6R. Cortes, C. Navarro, M.A. Martinez, J. Rodriguez and V. Sanchez-Galvez,
“Numerical modelling of normal impact on ceramic composite armors,” Int. J.
Impact Engng. 12[4], 639-651, (1992) 7 F. I. Grace, and N. L. Rupert, “Analysis of long rods impacting ceramic
targets at high velocity,” Int. J. Impact Engng. 20, 281-292 (1997),8 J.T. McGinn, R.W. Klopp and D.A Shockey, “Deformation and
comminution of Al2O3 in the Mescall zone of ceramic armor,” Ma.t Res. Soc.
Symp. Proc. 362, 61-66 (1995)9J.D. McGee, S. Nemat-Nasser, and J.B. Isaacs, “Ballistic performance of
ceramic tiles with thin membrane confinement,” submitted for publication10
S. Nemat-Nasser and J.Zhang, unpublished results 11
H. Horii and S. Nemat-Nasser, “Brittle failure in compression: Splitting,
faulting and brittle ductile transition,” Phil. Trans. Roy. Soc. Lond. 319[1549],
337-374 (1986) 12
S. Nemat-Nasser and H. Deng, “Strain-rate effect on brittle failure in
compression,” Acta Metall. Mater. 42[3], 1013-1024 (1994)
Ceramic Armor Materials by Design 525
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A NEW FAMILY OF REACTION BONDED CERAMICS FOR ARMOR
APPLICATIONS
M. K. Aghajanian, B. N. Morgan, J. R. Singh J. Mears and R. A. Wolffe
M Cubed Technologies, Inc. Simula Safety Systems, Inc.
1 Tralee Industrial Park 7822 South 46th
Street
Newark, DE 19711 Phoenix, AZ 85044
ABSTRACT
Reaction bonded SiC has existed for many years. It is produced by reactively
infiltrating a preform consisting of SiC and carbon with molten Si. During the
infiltration process, the Si and carbon react to form SiC. Thus, the finished body
consists of the original SiC, reaction-formed SiC, and residual Si. Historically,
reaction bonded SiC processes are designed such that high levels of reaction-
formed SiC are produced. With high levels of reaction formed SiC, the resultant
microstructure has a fully interconnected (coarse) SiC phase that provides good
performance in traditional ceramic applications (e.g., wear, corrosion, high
temperature, creep).
Within the past few years, new applications for ceramics have emerged in the
semiconductor industry (e.g., wafer chucks, wafer handling arms, process
chambers). These applications have different requirements than those for which
reaction bonded SiC was previously developed. For instance, the semiconductor
applications do not require wear or creep resistance, but do require excellent net
shape processing characteristics and a fine microstructure suited to machining of
minute details to high tolerance. To this end, a novel approach to reaction bonded
SiC was taken. Preforms that possessed a very high content of less than 50
micron SiC particles and a low carbon content were produced with a resin
molding process. Upon infiltration, little reaction occurred. This resulted in
minimal process shrinkage or distortion; and a microstructure with little
coarsening and low levels of residual stress. Such a material was found to be well
suited to near net shape production and machining to extremely precise tolerances
for semiconductor applications.
This novel reaction bonded SiC ceramic was evaluated for utility in armor
applications. The product was shown to possess good ballistic properties and the
ability to be produced to the desired tolerances without the need for machining.
Ceramic Armor Materials by Design 527
To the extent authorized under the laws of the United States of America, all copyright interests in this publication are the propertyof The American Ceramic Society. Any duplication, reproduction, or republication of this publication or any part thereof, withoutthe express written consent of The American Ceramic Society or fee paid to the Copyright Clearance Center, is prohibited.
Subsequently, a variation to the material was made. For applications requiring
greater hardness and lower density, a reaction bonded B4C ceramic was
formulated.
Herein, the processing, microstructure, properties, and ballistic performance
of these novel materials are presented and discussed.
INTRODUCTION
Reaction bonded SiC was first developed in the 1950’s 1,2,3
. Other terms for
the process include ‘reaction sintered’ and ‘self bonded’ 4. Conventionally, the
process consists of Si infiltration (liquid or vapor) into preforms of SiC + carbon.
During the infiltration step, the Si and carbon react to form SiC. Typically, all
carbon is consumed, yielding a product of porous SiC (vapor infiltration) or dense
Si/SiC (liquid infiltration). The maximum SiC particle size used in the production
of such bodies is generally in excess of a few hundred microns 1,2
.
A major advantage of the process is that the volume of the reaction-formed
SiC is 2.3 times larger than the volume of the reacted carbon. Thus, by
infiltrating Si into preforms that contain high carbon contents, ceramic bodies rich
in SiC can be produced.
Variations to the process have been studied. For example, Taylor and Palicka5
produced preforms of B4C and B4C + carbon and subsequently reactively
infiltrated the preforms with molten Si. The process conditions were selected to
encourage partial reaction between the Si and B4C, thus forming SiC (and
presumably SiBx). The resultant ceramic bodies contained B4C, SiC and Si. The
presence of B4C, which has a much lower density than SiC (2,540 vs. 3,210
kg/m3), yielded a ceramic body of low mass. To maximize the B4C content in the
components, a particle size distribution was utilized. A maximum B4C particle
size in the distribution of nominally 300 microns was chosen.
The reaction bonding process for SiC ceramics has several advantages relative
to traditional SiC processes (e.g., hot pressing, sintering). First and foremost,
volume change during processing is very low (generally well less than 1%),
which provides very good dimensional tolerance control. In addition, the process
requires relatively low process temperatures and no applied pressure, which
reduces capital and operating costs. Moreover, fine reactive powders capable of
being densified are not required, which reduces raw material cost. Finally, unlike
most monolithic SiC materials, Si/SiC is typically electrically conductive. This
allows EDM machining and assists in sensitive applications where static
discharge is required.
However, the vast majority of commercial reaction bonded SiC ceramics have
coarse microstructures. This is a due to the use of large SiC particles in the
preforms and the fact that many of these materials are made using high levels of
carbon in the preform. As the carbon reacts in an expansive manner with the Si to
528 Ceramic Armor Materials by Design
form SiC, the SiC particles in the preform are networked together to form large
SiC clusters. Since the strength of a ceramic is controlled by the largest flaw
within the stressed volume, a coarse grained material will tend to have low
strength. Therefore, reaction bonded SiC ceramics are traditionally used for high
temperature, creep, corrosion and wear sensitive applications, but not structural
(strength critical) applications.
The present work expanded on the aforementioned prior art with the goal of
producing optimized reaction bonded ceramic materials for room temperature
structural applications, such as semiconductor capital equipment components and
armor tiles. In particular, the activities focused on the production of components
with relatively fine-grained microstructures. Two different material types were
studied, namely reaction bonded SiC and reaction bonded B4C.
EXPERIMENTAL PROCEDURES
All of the reaction bonded ceramics described herein were produced with
nominally the same process steps. First, a preform was fabricated by mixing
ceramic particles with a resin binder, and casting the mixture into a mold. Next,
the mixture was cured, demolded, and exposed to about 600°C in an inert
atmosphere to pyrolyze the binder. Finally, the resultant carbon-bound preform
was contacted with a molten Si-containing alloy in a vacuum atmosphere, thus
allowing reactive infiltration to occur. Less than 0.5% volume change occurred
during the infiltration process.
After the fabrication step, various mechanical and physical properties of the
materials were measured. Density was determined by the water immersion
technique in accordance with ASTM Standard B 311. Elastic properties were
measured by an ultrasonic pulse echo technique following ASTM Standard D
2845. Hardness was measured on the Vicker’s scale with a 2 kg load per ASTM
Standard E 92. Flexural strength in four-point bending was determined following
MIL-STD-1942A. Fracture toughness was measured using a four-point-bend-
chevron-notch technique and a screw-driven Sintech model CITS-2000 universal
testing machine under displacement control at a crosshead speed of 1mm/min.
Specimens measuring 6 x 4.8 x 50 mm were tested with the loading direction
parallel to the 6 mm dimension and with inner and outer loading spans of 20 and
40 mm, respectively. The chevron notch, which was cut with a 0.3 mm wide
diamond blade, has an included angle of 60° and was located at the midlength of
each specimen. The dimensions of the specimen were chosen to minimize
analytical differences between two calculation methods according to the analyses
of Munz et al.6.
Microstructure was characterized in two manners. Polished sections were
examined using a Nikon Microphot-FX optical microscope. Fracture surfaces
were studied with a Jeol 840 scanning electron microscope (SEM).
Ceramic Armor Materials by Design 529
Advanced light armor designs typically consist of a ceramic tile to blunt or
break projectiles and a second layer (e.g., fiber-reinforced polymer composite) to
catch or stop the remains. For the present work, ballistic testing was conducted
using a simple configuration that simulated a typical light armor design.
Specifically, 100 mm x 100 mm ceramic tiles were bonded to 300 mm x 300 mm
fiber-reinforced polymer plates, and then were tested versus ballistic projectiles.
Ballistic resistance of the samples was determined by the procedures described in
MIL-STD-662F. Three materials were evaluated, namely reaction bonded SiC,
reaction bonded B4C, and commercial hot pressed B4C (control). In one series of
tests, the reaction bonded SiC and commercial hot pressed B4C were tested versus
ball rounds; and in a second set of tests the reaction bonded B4C and hot pressed
B4C were tested versus armor piercing (AP) rounds.
RESULTS AND DISCUSSION
Fabrication of Reaction Bonded SiC
The reaction bonded SiC ceramic material was produced in three basic steps.
First, a preform of SiC particles and organic resin was fabricated. Second, the
resin was pyrolyzed (converted to carbon). Finally, the preform was reactively
infiltrated with molten Si at nominally 1600°C. The final product was 100%
dense and consisted of the original SiC, reaction formed SiC (Si + carbon), and
remaining Si.
The goal of the reaction bonded SiC process development activities was to
produce a relatively fine grained ceramic for structural applications. To achieve
such a microstructure, the work utilized preforms with relatively small SiC
particles and low carbon content. The small particles led to a fine structure and
the low carbon content resulted in minimal reaction-formed SiC that would
cluster the small particles together. Specifically, a SiC particle size blend was
used to maximize particle packing. A maximum particle size of nominally 45
microns was used in the blend. The preforms produced with the blend contained
75 vol. % SiC and 4 vol. % carbon (pyrolyzed binder). After infiltration with
molten Si, the resultant bodies consisted of 84 vol. % SiC (75 original and 9
reaction formed) and 16 vol. % Si (i.e., an Si/SiC composite). A typical
microstructure (optical photomicrograph) of the material is shown in Figure 1.
In the optical photomicrograph, it is not possible to differentiate between the
original SiC and the reaction formed SiC. Nonetheless, it is clearly evident that
by the use of the relatively low carbon content little growth and interlocking of
the SiC particles has occurred, thus allowing a relatively fine microstructure to be
maintained.
530 Ceramic Armor Materials by Design
Fabrication of Reaction Bonded B4C
The reaction bonded B4C was produced in a nearly identical manner to the
reaction bonded SiC. A B4C particle blend with a maximum particle size of
nominally 45 microns was prepared. Preforms were then made using this blend.
The preforms consisted of nominally 75 vol. % B4C and 4 vol. % carbon. After
infiltration, the ceramic material contained nominally 75 vol. % B4C, 9 vol. %
reaction-formed SiC, and 16 vol. % remaining Si (i.e., an Si/SiC/B4C composite).
An optical photomicrograph of the material is shown in Figure 2.
Figure 1. Optical Photomicrograph of Reaction Bonded SiC
Figure 2. Optical Photomicrograph of Reaction Bonded B4C
As with the reaction bonded SiC, the reaction bonded B4C ceramic shown in
Figure 2 displays little interlocking and clustering of the particles. Also, the
photomicrograph shows little visible reaction between the Si and B4C as a result
Ceramic Armor Materials by Design 531
of the infiltration process. This was achieved by using process conditions
specifically designed to minimize reaction, including low process temperature,
short process time, and B-doping of the Si infiltrant. If Si-B4C reaction is allowed
to occur, as was the case in some previous work5, the microstructure significantly
coarsens. A coarse microstructure leads to a ceramic with a larger flaw size, and
thus lower strength. In Figure 3, a typical microstructure is shown were Si-B4C
reaction has occurred. Coarsening of the structure (i.e., large ceramic clusters
within the Si matrix) is clearly evident.
Figure 3. Optical Photomicrograph of Reaction Bonded B4C
with Unwanted Si-B4C Reaction
Mechanical and Physical Properties
Results of density, Young’s modulus, flexural strength and fracture toughness
are provided in Table I. When appropriate, the results are provided as a mean +/-
one standard deviation.
Table I. Properties of Reaction Bonded Ceramics
Property Reaction
Bonded SiC
Reaction
Bonded B4C
Density (kg/m3) 3060 2570
Young’s Modulus (GPa) 384 +/- 2 382 +/- 6
Flexural Strength (MPa) 284 +/- 14 278 +/- 14
Fracture Toughness (MPa-m1/2
) 3.9 +/- 0.5 5.0 +/- 0.4
The density of the SiC-based material is about 6% lower than monolithic SiC
due to the presence of the Si phase, which has relatively low density. This
reduced density is important for applications, such as armor, that are weight
532 Ceramic Armor Materials by Design
specific. The B4C-based material has very low density and is similar to that of
monolithic B4C.
The Young’s moduli of the reaction bonded SiC and reaction bonded B4C
ceramics are essentially the same, and compare favorably with other high
performance ceramic materials. The specific results are as predicted based on the
handbook Young’s modulus values for dense SiC, B4C and Si of ~450, ~450 and
120 GPa, respectively7. In particular, on a weight specific basis, the reaction
bonded B4C has a very high Young’s modulus.
The fracture toughness of the reaction bonded SiC of nominally 4 MPa-m1/2
,
is consistent with most SiC-based ceramics7. Surprisingly, the reaction bonded
B4C ceramic shows a 28% increase in toughness relative to the SiC material,
despite the fact that no ductile phase was added. A possible explanation for this
increased toughness was found by examining fracture surfaces, as is explained in
the next section.
Hardness is a very important parameter for armor materials. Previous work
has demonstrated that high mass efficiencies are only obtained versus hard armor
piercing projectiles when the projectiles are fractured, and that to effectively
fracture the projectile, an armor must have high hardness8,9
. However, it is
difficult to compare the many hardness data in the open literature because results
can be highly dependent on test method and technique. Therefore, for the present
work many different commercial materials were obtained. Hardness
measurements were then made on both the commercial materials and the new
reaction bonded ceramics in an identical manner so that true comparisons could
be made. The results are provided in Table II.
Table II. Results of Hardness Measurements
Material Vicker’s Hardness with
2 kg Load (kg/mm2)
7.62 mm M2 AP Bullet (Tool Steel) 926 +/- 26
14.5 mm BS-41 Bullet (WC/Co) 1644 +/- 30
Sintered AlN 1044 +/- 63
Pure Si 1243 +/- 21
90% Sintered Al2O3 1250 +/- 89
Hot Pressed AlN 1262 +/- 51
99.5% Sintered Al2O3 1499 +/- 74
Hot Pressed Al2O3 2057 +/- 82
Hot Pressed TiB2 2412 +/- 135
Hot Pressed TiC 2474 +/- 188
Hot Pressed SiC 2640 +/- 182
Hot Pressed B4C 3375 +/- 212
Ceramic Armor Materials by Design 533
Figure 4. SEM Fractographs of Reaction Bonded SiC (A)
and Reaction Bonded B4C (B)
Reaction Bonded SiC 2228 +/- 274
Reaction Bonded B4C 2807 +/- 54
The reaction bonded SiC and B4C ceramics have very high hardnesses that are
well in excess of both tool steel and WC/Co projectiles. In both cases, the Si/SiC
and Si/SiC/B4C composites have hardnesses that more-or-less reflect the weighted
average hardness of the constituents. In particular, because of the very high
hardness of monolithic B4C, the reaction bonded B4C has a very high hardness
value.
Analysis of Fracture Surfaces
The relatively high fracture toughness of the reaction bonded B4C ceramic
was unexpected. To gain an understanding for this result, the fracture surfaces of
the reaction bonded SiC and reaction bonded B4C ceramics were studied and
compared. The SEM fractographs for the two materials are provided in Figure 4.
534 Ceramic Armor Materials by Design
A significant difference between the two fracture surfaces is seen. The
reaction bonded SiC ceramic shows brittle, transgranular fracture of the SiC
particles. Also, brittle fracture of the Si matrix is seen. In addition some
indications of interfacial cracking between the Si and SiC are seen. The reaction
bonded B4C ceramic shows brittle, transgranular fracture of the B4C particles.
However, the Si matrix shows some highly unexpected ductile behavior with the
characteristic chisel-like rupture pattern. In addition, no evidence of failure at the
interfaces between the particles and matrix is seen. It is felt that the observed
semi-ductile failure of the Si phase is contributing to the relatively high toughness
of the reaction bonded B4C ceramic (Table I).
A review of the literature10-13
finds that Si undergoes a brittle to ductile
transition in the 500°C temperature range. The transition temperature decreases
as the dislocation density in the Si increases. In one study11
, more surface
dislocations were introduced to the surface of a sample by grinding, which
reduced the brittle to ductile transition temperature.
In the reaction bonded SiC system, little stress will be induced in the Si phase
on cooling from the processing temperature because both Si and SiC have CTEs
of nominally 4 ppm/K14
. Thus, the dislocation density in the Si should be low.
However, the situation is very different in the reaction bonded B4C ceramic.
Upon cooling from the process temperature, the B4C and Si will shrink at
different rates (B4C has a CTE of about 5.6 ppm/K14
). Thus, the Si will become
highly stressed and thus will have a high dislocation density. It is postulated that
this high dislocation density leads to the semi-ductile behavior of the Si in the
reaction bonded B4C ceramic at room temperature. More study of this
phenomenon is needed.
Ballistic Properties
The results of ballistic testing are provided in Tables III and IV. In Table III,
test results versus a 7.62 mm M80 ball round for reaction bonded SiC and
commercial hot pressed B4C (control) are provided. In Table IV, test results
versus a 7.62 mm AP M2 round for reaction bonded B4C and commercial hot
pressed B4C are provided. In each case, the tables provide the areal density of the
system, the mass efficiency of the target, and the normalized mass efficiency
relative to the hot pressed B4C control. The mass efficiencies in the tables were
determined based on available data for rolled homogeneous steel armor (RHA)
versus the same threats. Specifically the mass efficiency was calculated as the
areal density of RHA required to give the same performance divided by the areal
density of the tested targets.
Ceramic Armor Materials by Design 535
The ballistic results are very encouraging. They show that the armor designs
employing lower cost reaction bonded ceramics had mass efficiencies equivalent
to armors of the same design using hot pressed ceramics. This has enabled the
production of cost effective armor products for various applications, as is
discussed in the next section.
Table III. Ballistic Test Results versus 7.62 M80 Ball Threat
Armor System
Areal Density
kg/m2 (psf)
Mass Efficiency
(RHA Equivalent)
Normalized Mass
Efficiency
Hot Pressed B4C
(control)
23.5 (4.82) 4.56 1.00
Reaction Bonded
SiC
23.9 (4.89) 5.11 1.12
Table IV. Ballistic Test Results versus 7.62 AP M2 Threat
Armor System
Areal Density
kg/m2 (psf)
Mass Efficiency
(RHA Equivalent)
Normalized Mass
Efficiency
Hot Pressed B4C
(control)
29.0 (5.95) 4.53 1.00
Reaction Bonded
B4C
30.2 (6.18) 4.85 1.07
Examples of Products
Numerous armor products have been fabricated and tested using the novel
reaction bonded ceramic materials. Examples are provided in Figure 5. Key
process elements are that large components can be fabricated (no pressure
required), high tolerances can be obtained without machining (< 0.5% process
shrinkage), and costs are relatively low (no fine reactive powders, relatively low
fabrication temperatures). In Figure 5, the aircraft armor and personnel armor
tiles are fabricated from reaction bonded SiC. The vehicle armor plate is
fabricated from reaction bonded B4C.
Presently, reaction bonded SiC personnel armor plates are being manufactured
in very high volumes for the US Marine / US Army Interceptor program. In
addition, various aircraft and vehicle armor components are being produced in
lower volumes for both standard (reaction bonded SiC) and AP (reaction bonded
B4C) armor applications.
536 Ceramic Armor Materials by Design
Figure 5. Example Armor Products - Aircraft Armor Tiles (top), Vehicle Armor
Plate (middle), Personnel Armor Tiles (bottom).
Ceramic Armor Materials by Design 537
SUMMARY
Two new reaction bonded ceramics were developed, one based on SiC and
one based on B4C. In both cases, the process conditions were selected to yield a
fine-grained structure relative to traditional liquid-Si infiltrated reaction bonded
ceramics. Both materials show excellent mechanical properties and high
hardness. In particular, the reaction bonded B4C was found to have an
unexpectedly high fracture toughness. A proposed mechanism for the high
toughness was presented based on fracture surface analysis and previous
observations in the literature. Finally, the ballistic performance of the new
ceramics was measured. Relative to the incumbent hot pressed B4C, the reaction
bonded SiC showed good performance versus a 7.62 mm ball round and the
reaction bonded B4C showed good performance versus a 7.62 mm AP round.
REFERENCES 1 K.M. Taylor, “Cold Molded Dense Silicon Carbide Articles and Methods of
Making the Same,” U.S. Pat. No. 3 205 043, Sept. 7, 1965. 2 P.P. Popper, “Production of Dense Bodies of Silicon Carbide,” U.S. Pat. No.
3 275 722, Sept. 27, 1966. 3 C.W. Forrest, “Manufacture of Dense Bodies of Silicon Carbide,” U.S.
Patent No. 3 495 939, Feb. 17, 1970. 4 R. Morrell, Handbook of Properties of Technical and Engineering Ceramics,
HMSO Publications, London, U.K., 1985. 5 K.M. Taylor and R.J. Palicka, “Dense Carbide Composite for Armor and
Abrasives,” U.S. Pat. No. 3 765 300, Oct. 16, 1973. 6 D.G. Munz, J.L. Shannon, and R.T. Bubsey, “Fracture Toughness
Calculation from Maximum Load in Four Point Bend Tests of Chevron Notch
Specimens,” Int. J. Fracture, 16 R137-41 (1980). 7
Engineered Materials Handbook, Vol. 4, Ceramics and Glasses, ASM
International, Metals Park, Ohio, 1991. 8 M.L. Wilkins, R.L. Landingham, and C.A. Honodel, “Fifth Progress Report
of Light Armor Program,” Report No. UCRL-50980, University of CA,
Livermore, Jan. 1971. 9 C. Hsieh, “Ceramic-Faced Aluminum Armor Panel Development Studies,”
Appendix 9 of Report No. JPL-D-2092, Jet Propulsion Laboratory, Feb. 1985. 10
J. Samueles, S.G. Roberts, and P.B. Hirsch, “The Brittle-to-Ductile
Transition in Silicon,” Materials Science and Engineering, A105/106 39-46
(1988).11
P.D. Warren, “The Brittle-Ductile Transition in Silicon: The Influence of
Pre-Existing Dislocation Arrangements,” Scripta Met., 23 637-42 (1989). 12
K. Sumino, “Dislocations and Mechanical Properties of Silicon,” Materials
Science and Engineering, B4 335-41 (1989).
538 Ceramic Armor Materials by Design
13 P. Haasen, “Brittle-to-Ductile Transition in Silicon as a Model for
Intermetallics,” Materials Science and Engineering, A137 105-10 (1991). 14
Y.S. Touloukian [ed.], Thermophysical Properties of Matter, Plenum Press,
New York, 1970.
Ceramic Armor Materials by Design 539
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FLEXIBLE CERAMIC COATED FIBER FABRICS FOR LIGHT WEIGHT
PROTECTION SYSTEMS
Konstantin von Niessen and Rainer Gadow
University of Stuttgart
Institute for Manufacturing Technologies of Ceramic Components
and Composites (IMTCCC/IFKB)
Allmandring 7b
D-70569 Stuttgart, GERMANY
ABSTRACT
Based on thermal spray technologies a coating process for refractory oxide ce-
ramic layers even on temperature sensitive fiber substrates has been developed, so
that the coated fabrics retain their flexibility. High speed and high rate ceramic
coating is performed with simultaneous cooling so that refractory oxide ceramic
coatings can be applied on aramide and mullite fibers with potential for industrial
application. The penetration by bullets, knives and blades through such ceramic
coated multilayer fabrics is effectively prevented.
INTRODUCTION
For personnel protection as well as protection of aircrafts and cars, only light
and flexible materials can be used.1 Light and flexible fabrics made of aramide or
other high tenacity fibers meet some of these demands but their protection is not
sufficient. Sharp blades as well as high speed bullets can pierce these fabrics even
if several layers are used. This paper focuses on a new approach by coating fabrics
made of high tenacity fibers such as aramide and mullite fibers with a highly re-
fractory oxide ceramic by thermal spray technologies. By the combination of high
tenacity fiber woven fabrics and high performance ceramic coatings the penetra-
tion by bullets, knives and blades can be effectively prevented. The ceramic coat-
ing increases the fiber to fiber friction which prevents wave distortion and de-
lamination. The penetrating objects cannot change the fabric structure and push
the fibers aside. The hard oxide ceramic coating blunts sharp metal blades by
abrasion so they cannot trench the fabric, and the high friction between the ce-
ramic coating and the metal blade stops further penetration.
Ceramic Armor Materials by Design 541
To the extent authorized under the laws of the United States of America, all copyright interests in this publication are the propertyof The American Ceramic Society. Any duplication, reproduction, or republication of this publication or any part thereof, withoutthe express written consent of The American Ceramic Society or fee paid to the Copyright Clearance Center, is prohibited.
MATERIAL SCREENING
The material screening focusses on the use of high tenacity fiber fabrics and
highly refractory oxide ceramics. Two different commercially available fiber fab-
rics have been selected, the standard aramide fabric used for ballistic protection
Twaron CT 710 (Twaron Products, Wuppertal, Germany) and the mullite fiber
fabric Nextel 720 (3M, Minneapolis, MN, USA) consisting of 85% Al2O3 and
15% SiO2. The material properties of these fibers are summarized in table I.
Table I. Properties of fiber fabrics2
Fiber
fabricDensity
[g/cm3]
Tenacity
[MPa]
Initial
modulus E
[GPa]
De-
comp.
temp.
TD [°C]
Specific
heat CP
[J/kgK]
Max. appl.
Tem. TM
[°C]
Twaron
CT 710
1.45 2,800 85 500 1420 200
Nextel
720
3.40 2,100 260 2,000 800 1,204
Due to their high hardness and wear resistance the oxide ceramics Al2O3 and
TiO2 have been chosen as coating materials for thermal spraying. To improve the
bonding strength of the ceramic coatings on the fabric, AlSi is used as additional
bond coat. The bulk material properties of the ceramic materials are shown in ta-
ble II.
Table II. Bulk material properties of Al2O3 and TiO23
Oxide
ceramic
Density
[g/cm3]
Vickers
hardness
HV [-]
Youngs
modulus E
[GPa]
Melting
temp. TM
[°C]
Specific heat
CP [J/kgK]
Al2O3 3.98 2,200 400 2047 1,047
TiO2 4.25 1,150 205 1,860 730
In order to apply these oxide ceramics by thermal spraying, they have to be
available as spray powders. After a sintering process the used powders are me-
chanically broken and milled to a grain size of 10 – 22 µm.
542 Ceramic Armor Materials by Design
DEPOSITION OF OXIDE CERAMIC COATINGS ON LIGHTWEIGHT FIBER
FABRICS BY THERMAL SPRAYING
The thermal spray process allows the application of a broad variety of metal-
lurgical, cermetic and ceramic coatings on a variety of substrates. The Atmos-
pheric Plasma Spray (APS) process uses an electric arc discharge between a water
cooled copper anode and a tungsten cathode as an energy source. This electric arc
discharge dissociates and ionizes the working gas and builds up a plasma that ex-
pands into the atmosphere forming a plasma gas jet (see Fig.1).4
cooling
powder injection
plasma
cathode
(tungsten)
anode
plasmagas
���������������������
coating
substrate
courtesy Linde AG
energy source: el. arc / plasma
plasma temp.: up to 20.000 K
plasma gas: argon, helium, hydrogen, nitrogen
spray material: oxide ceramics, metals, alloys,
polymers
raw material form: powder
particle velocity: up to 450 m/s
deposition rate: 4 - 8 kg/h (oxide ceramics)
Fig. 1 The Atmospheric Plasma Spray (APS) process5
The spray powder, suspended in a carrier gas, is injected into the heat source of
the torch. After being totally or partially molten and being accelerated, the powder
particles impact on the substrate`s surface, where they are quenched and solidified
within 10-5
to 10-7
seconds. During atmospheric plasma spraying process tempera-
tures up to 20,000 °C are obtained.
For the application of thermally sprayed coatings on fiber woven fabrics the
torch movement is performed by a 6-axis robot system and a metal frame is used
for inserting and tentering the samples. The meandering movement and the metal
frame are shown in Fig. 2.
Two piece
metal frame
Screw joint
Wire cloth
Metal frame to
support and
stabilize
the fabric
APS plasmatorch
X
Y
Coating track configuration
Fig. 2 Mounting support for the fabrics and coating track configuration
In order to limit the thermal load on the fabrics a simultaneous cooling with
compressed air is used. Air nozzles are attached on both sides of the spraying
Ceramic Armor Materials by Design 543
torch. In addition, the process is supervised by an infrared camera (Varioscan In-
fraTec ID, Dresden, Germany) and in that way the temperature of the coated sam-
ples can be controlled in real time. Fig. 3 shows a typical IR- picture during the
coating process.
Fig. 3 IR- picture of the temp. distribution during the coating process
MECHANICAL CHARACTERIZATION
With regard to the use of the coatings for ballistic protection, the main focus of
the characterization is on the determination of puncture resistance, hardness and
wear resistance as well as on the evaluation of the coating`s bonding strength on
the first fiber layers. During the coating buildup of thermally sprayed layers, po-
rosity and microcracks cannot be avoided. For the coating of flexible fabrics the
formation of porosity and microcracks in the coating is desired because it leads to
a higher flexibility of the fabric. But if the porosity is too high, the hardness and
other mechanical properties of thermally sprayed coatings decrease. So a balance
between porosity and mechanical properties has to be found.
The thickness of the oxide ceramic coatings on the fabric is in the range of 50 –
100 µm. Fig 4 shows a schematic drawing of the intended structure of the coated
fabric.
544 Ceramic Armor Materials by Design
50-100µm
Fiber woven
fabric
Fig. 4 Intended structure of the oxide ceramic coated fabric
In Fig. 5 a cross section of a Twaron fabric coated with an Al2O3 oxide ce-
ramic layer is shown. The lamellar structure and the good wetting behavior of the
ceramic coating on the first layers of the fabric are visible. The macro-structure
and micro-structure of the coated fabric`s surface is typical for thermally sprayed
coatings (see Fig. 6). The structure of the fabric is still visible in the macro-
structure. Even though the TiO2- and Al2O3- coatings have melting points above
1800° and 2000°C respectively, there is no significant polymer fiber damage.
Fig. 5 Cross section of a thermally sprayed Al2O3 coating on a Twaron fabric
Al2O3- Coating
Twaron Fabric
Ceramic Armor Materials by Design 545
Fig. 6 SEM micrographs of a thermally sprayed Al2O3 coating on a Twaron
fabric
Macro structure Micro structure
In order to evaluate the coating quality metallographic examinations have been
performed. The coating porosity determined by an image analysis is expressed by
the relative pore volume content VP [%]. An automized universal hardness in-
denter equipment (Fischerscope TM HCU) with a load of 500 mN is used to de-
termine the coating hardness HV0,05. In order to measure the hardness of an indi-
vidual fiber, the load was reduced to 10 mN (HV 0,001). Table III and table IV
show the measured porosity and hardness characteristics of the thermally sprayed
coatings and of the fibers, respectively.
Table III. Measured coating porosity and hardness (HV 0,05)
Coating VP [%] HV 0,05
Al2O3 5.8 1,240 +/- 300
TiO2 3.2 1,100 +/- 110
Al2O3/TiO2 4.1 1,025 +/- 180
AlSi 1.44 138 +/- 10
Table IV Microhardness of individual fibers (HV 0,001)
Fiber HV 0,001
Twaron CT 710 51.52 +/- 7
Nextel 720 1,610 +/- 405
The investigation of the coating`s adhesion on the fabric is performed on a
Zwick Z100 universal mechanical testing machine by pull testing. The coated fab-
ric samples are glued to a metal plate and a steel tension rod is glued to the coat-
ing surface by using an adhesive. After mounting the samples into the testing ma-
chine the tensile load is continuously increased. As soon as a delamination of the
coating occurs, the tensile load is measured and the bonding strength is deter-
546 Ceramic Armor Materials by Design
mined. As the bonding strength of the coatings is limited by the maximum shear
strength of the first fiber layers which are in contact with the coating, the fabrics
are also tested without any coating. In this case the tensile rod is glued directly on
top of the fabric. Fig. 7 shows the measured bonding strengths for the used fabrics
with or without AlSi bond coat.
without coatin
Al 2O 3
TiO
2
Al 2O
3/T
iO2
AlS
i
Al 2O
3 - A
lSi
TiO
2 - A
lSi
Al 2O
3/T
iO2 - A
lSi
Nextel 720
Twaron
0
1
2
3
4
5
6
7
Fig. 7 Bonding strengths of the thermally sprayed coatings on fiber fabrics
Bo
nd
ing
str
en
gth
[N
/mm
2]
The results of the experiments with non–coated fabrics show the maximum
possible bonding strength a coating could reach on the fabrics. Because of its low
shear strength, Nextel already reaches its limit at a bonding strength of 3 N/mm2.
The TiO2 coatings reach this values with and without a AlSi bond coat. The bond-
ing strength of the Al2O3 and Al2O3/TiO2 coatings is rather low, however it can be
increased by using the additional AlSi bond coat. For the TiO2 coated Nextel
fabrics delamination occurs within the fabric itself, whereas the other coatings
with lower bonding strength delaminate at the fiber–coating interface. Due to a
higher shear strength the maximum bonding strength of Twaron is about 7
N/mm2. None of the oxide ceramic coatings reach this limit, but by the use of a
bond coat, the bonding strength on Twaron is increased. Especially the Al2O3-
AlSi coating shows a high bonding strength and the highest microhardness. All
coatings deposited on the Twaron fabrics delaminated at the fiber–coating inter-
face. The differences in the mechanical properties of the coated fabrics are obvi-
Ceramic Armor Materials by Design 547
ous. This might be due to the differences in the process temperatures since the
fabrics have different thermophysical properties, which influence the wetting and
bonding behavior of the applied coatings.
Because of the good results obtained with the Al2O3-AlSi coating on Twaron
fabrics, comparative stab resistance tests on these coatings and uncoated Twaron
fabrics are performed. German standard engineered test blades K1 (A. Eickhorn
GmbH, Solingen, Germany) for stab resistance tests are mounted into a Zwick
Z100 universal mechanical testing machine. The fabrics are fixed in a specific de-
vice by hydraulic pressure to obtain a defined prestress. The puncture resistance
performance is measured in work [N mm] per penetration depth [mm]. In one
experimental run, 6 stabs are carried out on different samples of the same fabric,
using one test blade to evaluate the blunting of the blades. The test velocity of the
blade is varied from 50 to 1500 mm/min, but no influence on the results was ob-
served. The typical run of the curves show an increase of the puncture resistance
for every new stab. This increase, which is caused by the blunting of the blade, is
for the Al2O3-AlSi coated Twaron fabrics much higher than for the uncoated fab-
rics. The penetration work of the coated fabrics is 5 times higher in comparison to
the uncoated fabrics. Fig. 8 shows the measured puncture resistance of Al2O3-AlSi
coated Twaron fabrics and uncoated Twaron fabrics.
0 20 40 60
Fig. 8 Stab resistence of Al2O3-AlSi coated and uncoated Twaron fabric
80
0
500
1000
1500
2000
Penetration depht [mm]
Pen
etra
tion w
ork
[Nm
m]
Stab resistance in N*mm of Twaron
fabric coated with Al3O3-AlSi multi-
layer coating
Stab resistance in N*mm of uncoated
Twaron fabric
Twaron CT 710
1 layer
plain weave style
Increase of the penetration work for
every new stab, which is caused by
the blunting of the metal blade
548 Ceramic Armor Materials by Design
CONCLUSIONS
The approach to combine highly refractory oxide ceramic coatings with high
modulus lightweight fiber fabrics has been successfully demonstrated. Atmos-
pheric plasma spraying with well defined parameter sets and simultaneous cooling
is a suitable process for the coating of oxide ceramics on top of fiber woven fab-
rics for ballistic protection. Even though the TiO2- and Al2O3- coatings have melt-
ing points above 1800° and. 2000°C respectively, no significant polymer fiber
damage has been seen. The adherent coatings remain flexible and reach a hardness
up to 1240 HV 0,05. The bonding strength is sufficient and mainly limited by the
maximum shear strength of the fibers. The adhesion of the coatings and the high
cycle flexibility can be improved by using metallurgical bond coats. So far the
best results have been reached with an Al2O3- coating on a Twaron fabric with a
AlSi bond coat. It has the highest microhardness and the highest bonding strength.
Stab resistance tests were carried out on Al2O3-AlSi multilayer coated Twaron
fabrics and the penetration work was increased by a factor of five compared to the
uncoated Twaron fabric. Further efforts will focus on the optimization of the in-
terface between oxide ceramic coating and fiber fabrics by tailoring the cooling
process during thermal spraying as well as by deposition of metallurgical thin
films as bond coats.
ACKNOWLEDGMENT
The authors would like to thank Mrs Katrin Keck (metallography) and Mr
Chuanfei Li (plasma spraying) for their support and Mr Scherer for the helpful
discussions.
REFERENCES 1 J-P. Charles, D. Guedra- Degeorges, “Impact Damage Tolerance of Helicop-
ter Sandwich Structures,” Aerospatiale, France (1999) 2“Product data sheet Twaron”, Twaro Products, D- 42097 Wuppertal, Kasi-
nostraße 19-21, (1995) 3 C. Friedrich, G. Berg, E. Broszeit, C. Berger: “Datensammlung zu Hartstof-
feigenschaften,” Materialwissenschaft und Werkstofftechnik, Vol. 28, No. 2,
(1997)4 L. Pawlowski, “The science and engineering of thermal spray coatings,” John
Wiley and sons, Chichester (1995)5“Das Verfahrensspektrum beim thermischen Spritzen,“ Linde AG, Werks-
gruppe technische Gase, Höffriegelskreuth (1990)
Ceramic Armor Materials by Design 549
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IMPROVED PERFORMANCE OF ALUMINA CERAMICS WITH CARBON
NANOTUBE REINFORCEMENT
Michael Sennett
U.S. Army SBCCOM
Natick Soldier Center
Natick, MA 01776-5020
Sekyung Chang, Robert H. Doremus, Richard W. Siegel, Pulickel M. Ajayan and
Linda S. Schadler
Materials Science and Engineering Department and
Rensselaer Nanotechnology Center
Rensselaer Polytechnic Institute
Troy, NY 12180-3590
ABSTRACT
Nanoscale alumina powder and carbon nanotubes were mixed and hot-pressed
to form dense ceramic-matrix composites. The strength and fracture toughness of
hot-pressed alpha-alumina was much greater than that of conventional grain size
polycrystalline alumina. The addition of carbon nanotubes to the alumina resulted in
composites with even greater strength and fracture toughness. Hot pressing in a
vacuum improved both of these properties over hot pressing in argon. These results
suggest that lightweight composites of high strength and fracture toughness can be
made from composites of nanophase alumina, or other ceramics, and carbon
nanotubes.
INTRODUCTION
Carbon nanotubes have high modulus and aspect ratio1,2
, and thus may be
excellent reinforcing fillers for ceramics. The mechanical properties of such
composites will depend strongly on the processing method and surface treatment of
the carbon nanotubes. Sintered alumina has high strength, hardness, and fracture
toughness. Improving these properties by incorporating carbon nanotubes in an
alumina-matrix composite is an exciting possibility as well as a processing challenge.
Here we report on the processing and mechanical properties of composites
made from nanoscale alumina particles to form the matrix and multi-wall carbon
Ceramic Armor Materials by Design 551
To the extent authorized under the laws of the United States of America, all copyright interests in this publication are the propertyof The American Ceramic Society. Any duplication, reproduction, or republication of this publication or any part thereof, withoutthe express written consent of The American Ceramic Society or fee paid to the Copyright Clearance Center, is prohibited.
nanotubes (MWNT) as the reinforcing material. We give special emphasis to
improved methods of dispersing the MWNT in the alumina powder before pressing
and sintering, and to purification and oxidation of the MWNT.
EXPERIMENTAL PROCEDURE
Gamma-phase alumina powder consisting of particles with a mean diameter
of 23 nm (Nanophase Technologies Corporation, Romeoville, IL) and MWNT,
synthesized by the arc-discharge method, were used to make the composites. The
gamma-phase alumina powder was transformed to alpha-alumina before sintering by
heating at 1300oC for 7 min. The mean particle size of the alpha powder determined
from X-ray line broadening was about 62 nm.
Alumina matrix composites with 5-20 vol.% MWNT were fabricated. The
MWNT were lightly oxidized by heating them at 640oC in air for various lengths of
time up to 150 min. This treatment removes some of the carbonaceous material and
makes it easier to disperse the nanotubes. The alpha-alumina powder and MWNT
were dispersed in dichloromethane (methylene chloride, CH2Cl2) with an ultrasonic
probe for about 4 min. The mixture of alumina and MWNT was held in the
ultrasonic bath until most of the CH2Cl2 evaporated, and then the mixture was dried
at 75oC for 24 hr. The weakly agglomerated mixture was ground and remixed in an
agate mortar and pestle and then dried at 130oC for 12 hr. Finally, the alumina-
MWNT mixtures were sintered by hot pressing in a graphite die at 1300oC, and a
pressure of 60 MPa, for 1 hr in an Ar atmosphere or in a vacuum hot press.
Alumina with un-oxidized MWNT composites, marked as “as received” were
prepared in the same way described in Reference 4
The density of the composites was measured by the Archimedes method. X-
ray analysis was performed on the composites to determine if the presence of the
MWNT causes the formation of any new phases.
To measure the hardness and fracture toughness, the surface of the
composites was polished with 1 µm diamond paste and then 0.3 µm alumina powder.
The hardness of the composites was measured with a micro-Vickers hardness
indenter (Model M-400, Leco Co.); a 1 kg load was applied on the surface for 10 sec.
To measure the fracture toughness of the composites, a Vickers hardness tester
(Vickers Limited) with a load of 5 kg was used, and the fracture toughness was
calculated from the lengths of cracks emanating from the indenter corners by the
“Evans & Charles” equation (Kc = 0.00824*(P/C1,5
), where P is equal to the applied
load in Newtons and C is equal to the crack length in meters.
The strength of the composite samples was measured with diametral tests of
sintered discs. In these tests a compressive load P is applied across the diameter d of
a disc sample. The result is a line of tensile stress along the sample surface and
through its volume to the other surface5,6
. The maximum stress S, which occurs on
552 Ceramic Armor Materials by Design
the diametral plane between loading points, is S = 2P/dL, in which d is the diameter
of the disc (16 mm in these tests) and L its thickness (4 mm). A pad of soft material
(copper) is inserted between the hard loading plates and the specimen.
Vardar and Finnie5 compared the strength of “granodiorite” (presumably
grandidierite, (Mg, Fe) Al3BSiO9) and limestone measured in bending (tensile) and
diametral tests over a wide range of strengths. They found quite similar Weibull
distributions and mean strengths for both of these minerals in the two tests.
Grandidierite has a hardness of 7.5 and limestone is of course soft, so these results
demonstrate the validity of the diametral test as compared to bending tests over a
wide range of hardness and strength.
Polished surfaces of composites and hardness indents were examined with
optical microscopy. Fracture surfaces were coated with gold and examined in a
scanning electron microscope (SEM, JEOL-A 40).
RESULTS
The samples described in this section were prepared by hot pressing in Ar at
1300oC and 60 MPa for 1 h unless otherwise noted. X-ray diffraction patterns
showed that the composites consisted of alpha-alumina and graphitic carbon only.
Broadening of the graphite diffraction lines showed that the average diameter of the
MWNT was about 12 nm (see also Ref. 4). The structures of the MWNT were the
same before and after processing. The density of the sintered composites was above
97% of theoretical density.
The diametral strengths of alumina-MWNT composites with different
MWNT content are shown in Fig. 1. For each MWNT content three as-received
samples were tested and the mean taken; one sample of each specimen with MWNT
dispersed in CH2Cl2 was fractured. The bulk alumina made from nanoparticles alone
had a strength of 654 MPa, which is greater than the typical strength of 200 to 350
MPa for sintered alumina7
and was even comparable with strengths reported for
single-crystal alumina (sapphire) of from 350 to 1000 MPa. The composites
containing as-received MWNT had somewhat lower strengths than that of the bulk
alumina. When carbon nanotubes dispersed in CH2Cl2 were added to alumina to
form composites, the strength first increased at 5 and 10 vol.% MWNT and then
decreased to the strength of composites with as-received MWNT at 20 vol.%
MWNT.
The fracture toughness of bulk alumina and alumina-MWNT composites is
shown in Fig. 2. The average fracture toughness of bulk alumina and alumina-5
vol.% MWNT composites, hot-pressed in a vacuum, increased to about 4.9 MPa m
and 5.1 MPa m, respectively. These toughness values are higher than those reported
for single-crystal alumina (sapphire) and polycrystalline alumina.8
Ceramic Armor Materials by Design 553
0 5 10 15 20 250
200
400
600
800
0 5 10 15 20 250
200
400
600
800
Content of MWNT ( vol. % )
MWNT ( as received )
MWNT ( purified in CH2Cl
2 )
Str
en
gth
( M
N/m
2 )
Fig. 1. Diametral strengths of alumina-matrix composites hot-pressed in Ar at
1300oC and 60 MPa for 1 h containing different amounts of MWNT.
Fig. 2. Fracture toughness of sintered nanophase alumina and alumina-MWNT
composites: black triangles, hot-pressed in argon; open triangles, vacuum hot
pressed.
0 5 10 15 20 250
1
2
3
4
5
6
Content of MW NT( vol. % )
MW NT ( Purified in CH2Cl
2 )
KC (
MP
a.m
0.5 )
MW NT ( as received )
554 Ceramic Armor Materials by Design
In previous work,4
the Vickers hardness of alumina composites containing as-
received MWNT decreased linearly from the bulk pure alumina value of 18.4 GPa to
13.5 GPa at 20 vol.% MWNT. In the present work, the hardness of an alumina -10
vol.% MWNT composite increased from 16.2 GPa with no oxidation of the MWNT
to a maximum hardness of 20.4 GPa after 90 min. of heating the MWNT in air at
640 C.
SEMs from the fracture surfaces of alumina-MWNT composites are shown in
Fig. 3. They show that the MWNT purified in dichloromethane are more evenly
dispersed than in composites made from as-received MWNT.
Fig. 3. Scanning electron micrographs of fracture surfaces of nanophase alumina-5
vol.% MWNT composites (a) and (b), as-received MWNT, (c) and (d) MWNT
purified in dichloromethane, all hot-pressed in argon.
(b)
(c) (d)
(a)
1 m3 m
3 m3 m
(b)
(c) (d)
(a)
1 m3 m
3 m3 m
DISCUSSION
The strength and fracture toughness of bulk alumina hot-pressed from
Ceramic Armor Materials by Design 555
nanophase alumina powder in Ar was much higher than typical strength and
toughness of the conventional polycrystalline alumina. The addition of 5 vol.%
MWNT to nanophase alumina to form a composite increased both the diametral
strength and fracture toughness even more. Purification of the MWNT in
dichloromethane improved the dispersion of the MWNT in the final composites. This
purification step removes excess carbon from the MWNT samples, leaving purer
MWNT that disperse better in the dichloromethane solvent. Vacuum hot pressing
removes entrapped gases in the composite powder mixture, preventing formation of
residual stresses and reduction of strength. These processing improvements show the
great promise of nanophase alumina-MWNT composites for lightweight, high-
strength materials. Further improvements in processing should lead to composites
with substantial content of MWNT and consequent low density, and having high
strength and fracture toughness.
ACKNOWLEDGEMENTS
This work was supported by the U.S. Army SBCCOM, Natick Soldier Center. We
thank Nanophase Technologies Corporation for supplying the nanophase alumina.
REFERENCES 1S. Iijima, “Helical Microtubules of Graphitic Carbon,” Nature, 35 [7] Nov.
56-58 (1991). 2O. Lourie and H. D. Wagner, “Evaluation of Young’s Modulus of Carbon
Nanotubes by Micro-Raman Spectroscopy,” J. Mater. Res., 13[9] 2418-2422 (1988). 3P. M. Ajayan and T. W. Ebbesen, “Nanometre-size Tubes of Carbon,” Rep.
Prog. Phys., 60 1025-1062 (1999). 4S. Chang, R. H. Doremus, P. M. Ajayan and R. W. Siegel, “Processing and
Mechanical Properties of C-Nanotube Reinforced Alumina Composites, Ceramic
Engineering and Science Proceedings, 21[3] 653-658 (2000). 5O. Vardar and I. Finnie, “An Analysis of the Brazilian Disk Fracture Test
Using the Weibull Probabistic Treatment of Brittle Strength,” Int. J. Fracture, 11 [3]
495-508 (1975). 6M. B. Thomas, R. H. Doremus, M. Jarcho and R. L. Salsbury, “Dense
Hydroxylapatite: Fatigue and Fracture Strength after Various Treatments, from
Diametral Tests,” J. Mat. Sci., 15 891-896 (1980). 7W. D. Kingery, H. K. Bowen and D. R. Uhlmann, Introduction to Ceramics,
John Wiley and Co., New York, 1976, p. 791. 8Y.-M. Chiang, D. P. Birnie, and W. D. Kingery, Physical Ceramics, John
Wiley and Co., New York, 1997, p. 484.
556 Ceramic Armor Materials by Design
RECENT PROGRESS ON THE INFLUENCE OF MICROSTRUCTURE AND
MECHANICAL PROPERTIES ON BALLISTIC PERFORMANCE
J.C. LaSalvia
U.S. Army Research Laboratory
Aberdeen Proving Ground, MD 21005-5069
ABSTRACT
Recent work on a terminal ballistic phenomenon known as dwell has led to
the identification of the important ceramic characteristics that govern this
phenomenon. In the ballistics community, dwell is used to describe the non-
penetration phase (complete or partial) of a long-rod projectile impacting on a
target. Because of the typical densities and velocities of long-rod projectiles,
dwell is typically observed in targets containing ceramics with high hardness
values. Recovery of ceramics from experiments in which complete dwell
occurred has led to the observation and basic understanding of the damage
mechanisms. Most notable is the importance of shear with respect to these
mechanisms. Consequently, a model for the transition from dwell-to-penetration
(a ballistic performance measure) was formulated by combining a
micromechanics-based compressive failure model with Hertz’s theory for
frictionless contact between axisymmetric linear-elastic bodies. The resulting
model indicates the relative importance of a ceramic’s grain size, short-crack
fracture toughness, yield strength, Poisson’s ratio, coefficient of friction, and
critical crack-length on the dwell/penetration transition. The motivation,
derivation, and predictions of the model are presented.
INTRODUCTION
Despite over 30 years of research and development of ceramic-based armor
technologies1,2
, a coherent and comprehensive understanding of the effect of a
ceramic’s physical and mechanical characteristics on its ballistic performance
does not exist. While there have been a few notable attempts at identifying and
correlating the important physical and mechanical attributes of a ceramic with
performance3,4
, these remain qualitative and do not allow performance to be
predicted. Another major problem has been the reliance of ballistic performance
measure methodologies (e.g. depth-of-penetration, V50) that do not adequately
Ceramic Armor Materials by Design 557
To the extent authorized under the laws of the United States of America, all copyright interests in this publication are the propertyof The American Ceramic Society. Any duplication, reproduction, or republication of this publication or any part thereof, withoutthe express written consent of The American Ceramic Society or fee paid to the Copyright Clearance Center, is prohibited.
(a) (b)
Figure 1. (a) Early confinement scheme used by Hauver et al.9. (b) X-ray flash
radiograph of a impacting long-rod dwelling on the surface of the ceramic
(normal impact).
distinguish the effect of a ceramic’s characteristics from the total system
performance nor yield sufficient insight into the fundamental processes that lead
to ceramic failure (i.e. penetration)5,6
.
Recently, tremendous insight into the fundamental projectile/target interaction
was gained through the work by Hauver et al.7-9
, Lundberg et al.10,11
, and others12-
16. Hauver et al.
7-9 discovered that through proper target design, the projectile
could be completely defeated without penetrating the ceramic. As a result of this
discovery, two new terms were added to the terminal ballistics vocabulary,
“dwell” and “interface-defeat”. Dwell is used to describe the state of the
projectile/target interaction event where the projectile does not penetrate the
ceramic and therefore has a zero penetration velocity. Interface defeat is used to
describe the condition when there is no significant penetration of the ceramic by
projectile during the entire ballistic event.
Lundberg et al.10,11
followed Hauver’s work with several fundamental studies
on not only the effect of projectile velocity on the penetration velocity through the
ceramic, but also on the projectile velocity where the onset of penetration
occurred (i.e. below this velocity, complete dwell occurred). This work coupled
with observations made on recovered ceramics from successful interface defeat
experiments and application of the results from a micromechanics-based
compressive failure model by Shih17
led LaSalvia et al.18
to develop a physically-
based theory that provided a rationale explanation for both the localized damage
and dwell/penetration transition velocity observations.
558 Ceramic Armor Materials by Design
5 mm5 mm5 mm5 mm5 mm5 mm5 mm5 mm5 mm5 mm5 mm
Figure 2. Cross-section of a titanium diboride tile recovered following a
successful interface defeat experiment17
.
The purpose of this paper is to provide a brief review of not only the theory
developed by LaSalvia et al.18
, but also the foundational work that led to it’s
development. A figure-of-merit is derived and its implications for connecting
microstructure and mechanical properties with the dwell/penetration transition
velocity are also presented.
BACKGROUND
Dwell and Interface Defeat
Dwell and interface defeat was first reported by Hauver and Melani7 in small-
scale reverse-ballistic experiments with heavily confined ceramic targets.
Subsequently, Hauver et al.9 conducted larger scale forward-ballistic experiments
in which subscale long-rod projectiles (L/D = 10 and 20, D = 5 - 6 mm) were
launched into ceramic targets at velocities up to 2000 m/s. The ceramics were
nominally 75 mm in diameter and 25 mm thick. Figure 1(a) is a schematic
illustration of an early confinement scheme that was used. In addition to the
ceramic being heavily confined on all sides, a shock-wave attenuator and a
tailored ceramic/front steel cover plate interface were incorporated in these larger
scale experiments. An X-ray radiograph taken during a dwell experiment is
shown in Figure 1(b)9. As can be seen, the long-rod projectile is dwelling on the
ceramic. Using the confinement scheme shown in Figure 1(a), Hauver et al.9
were able to achieve interface defeat against long-rod projectiles impacting at
1600 m/s for silicon carbide, titanium diboride, titanium carbide, and tungsten
carbide.
An important aspect of the larger-scale experiments conducted by Hauver et
al.9 was that the ceramics could be recovered and examined after ceramographic
preparation. Figure 2 show a cross-section from a recovered titanium diboride tile
that was impacted at 1600 m/s. In addition to the numerous lateral and cone
Ceramic Armor Materials by Design 559
cracks, a region of severe localized damage just beneath where the long-rod
impacted is clearly evident in the titanium diboride tile. As can be seen, this
localized damage region (often referred to as the “comminuted” region) does not
extend to the top surface, but is apparently fully confined by “undamaged”
material. The shape of the comminuted region corresponds well with calculated
and/or observed deviatoric stress distributions in quasi-static and dynamic contact
problems19-21
. This indicates the importance of shear with respect to the damage
mechanisms. With the exception of tungsten carbide, those ceramics that were
recovered after successful interface defeat experiments exhibited this localized
damage region.
Dwell/Penetration Transition Velocity
Although Hauver and Melani7 had discovered the dwell phenomenon and
interface defeat, much of our fundamental understanding is due to the
experimental work of Lundberg et al.10,11
. Using the reverse-ballistic testing
methodology, Lundberg et al.10,11
systematically investigated both the
dwell/penetration transition and the penetration rate for a number of ceramics.
Highly confined (see Figure 3(a)) boron carbide, silicon carbide, Syndie+, and
titanium diboride specimens were launched into either tungsten- and
molybdenum-based subscale long-rods (L/D = 40, D = 2mm) at velocities up to
2500 m/s. The dwell/penetration transition velocities for these ceramics against
tungsten-based rods are plotted in Figure 3(b) as a function of their estimated
compressive yield strengths ( Y). The yield strengths for the ceramics were
calculated from hardness measurements22-24
. As can be seen, the data for the
ceramics evaluated fall nominally on two curves; hereafter referred to as the upper
and lower bound curves. Silicon carbide and titanium diboride fall on the upper
bound curve, while boron carbide and Syndie fall on the lower bound curve.
Lundberg et al.11
offered the following explanations for these two curves. The
upper bound curve corresponds to the critical pressure required to form an indent
on a rigid, perfectly-plastic body using a rigid punch. The lower bound curve
corresponds to the critical pressure required to initiate yielding beneath the area
loaded. At first thought, both explanations appear to be highly questionable,
given the fact that ceramics do not typically have a sufficient number of
independent slip systems to support any appreciable amount of bulk plasticity in
response to loads. However, with sufficient confinement, bulk plasticity is
possible as is known in the compressive failure of geological materials25-30
. If a
ceramic is physically confined as shown in Figures 1(a) and 3(a), is ductile failure
of the ceramic possible? The data in Figure 3(b) supports the possibility for
ductile failure (i.e. the upper curve). However, the hypothesis of a completely
+De Beers, Inc. synthetic polycrystalline diamond.
560 Ceramic Armor Materials by Design
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(a) (b)
Figure 3. (a) Schematic illustration of the confined ceramic design used by
Lundberg et al.11
. (b) Dwell/penetration transition velocity data for several
ceramics plotted as a function of the ceramic’s yield strength11
.
ductile response (i.e. explanation for both the upper and lower curves) is not
consistent with the damage observations from ceramics recovered by Hauver et
al.9. In order to reconcile the data shown in Figure 3(b) with the observed damage
in the recovered ceramics, the compressive failure of brittle solids must first be
considered.
Compressive Failure of Brittle Solids
The general features of compressive failure of geological materials has been
the subject of a large number of investigations given its’ importance in the proper
laying of foundations for concrete structures25-30
. The effect of confining pressure
on the observed failure mechanism (e.g. axial splitting, faulting, and plastic flow)
was identified. This led to a number of compression failure models that are based
upon growth or suppression of microcracks with the wing-crack or z-crack flaw
geometry28-30
. Wing-cracks are mixed-mode cracks that extend out from the
plane of the pre-existing flaw, nominally in the direction of the maximum
principal stress. A schematic illustration of a wing-crack is shown in Figure 4(b).
In general, the initiation and initial growth of the wing-cracks are governed by the
local shear failure (mode II) of the pre-existing flaw and the ability of the
surrounding material to accommodate this failure by plastic deformation. As
shown in Figure 4, this shear failure can be accommodated either by plastic
deformation, wing-crack formation, or a combination of both. The ability to
accommodate the shear failure by microplasticity is an important material
Ceramic Armor Materials by Design 561
2
1
2
1
2
1
2
1
2
1
2
1
2
1
2
1
2
1
2
1
2
1
(a) (b) (c)
Figure 4. (a) Pre-existing flaw subjected to localized normal and shear stresses
due to far-field principal stresses 1 and 2. (b) Wing-crack initiation and growth
due to shear failure of pre-existing flaw. (c) Accommodation of shear failure of
the pre-existing flaw by dislocation generation.
characteristic since it leads to a suppression of wing-crack formation, and hence a
decreased potential for brittle failure (macro).
According to the compressive failure model proposed by Horii and Nemat-
Nasser29
, the propensity of a material to suppress the initiation and growth of
wing-cracks is indicated, though not exclusively, by its ductility parameter. The
ductility parameter is defined as29
:
2
cK
Y
IC* (1)
where KIC is the mode I fracture toughness, Y is the uniaxial compressive yield
strength for ductile failure, and 2c is the pre-existing flaw size. The ductility
parameter is dimensionless and represents the ratio of a brittle failure strength
measure to a ductile failure strength measure. A low ductility parameter would
indicate a stronger tendency towards wing-crack initiation and growth, while a
high ductility parameter would indicate a stronger tendency towards wing-crack
suppression.
The compressive failure model that was developed by Horii and Nemat-
Nasser29
is shown in Figure 5. A pre-existing flaw of length 2c is subjected to
far-field principal stresses 1 and 2. The pre-existing flaw makes an angle as
measured from the maximum principal stress direction. A plastic zone of length
and wing-crack of length are possible as a result of the sliding motion
(shear) of the pre-existing flaw surfaces. In the model, the plastic zones are
represented as collinear arrays of edge dislocations, while the wing-cracks extend
out at an angle . A frictional force, due to the combined effects of the resolved
pλ tλ
562 Ceramic Armor Materials by Design
2
1
(2c)Pre-Existing Flawpλ
tλ
2
1
(2c)Pre-Existing Flawpλ
tλ
2
1
(2c)Pre-Existing Flaw
2
1
(2c)Pre-Existing Flaw
2
1
(2c)Pre-Existing Flaw
(2c)Pre-Existing Flawpλ
tλ
2
1
(2c)Pre-Existing Flawpλ
tλ
2
1
(2c)Pre-Existing Flaw
2
1
(2c)Pre-Existing Flaw
2
1
(2c)Pre-Existing Flaw
(2c)Pre-Existing Flawpλpλ
tλtλ
2
1
(2c)Pre-Existing Flaw
2
1
(2c)Pre-Existing Flaw
2
1
(2c)Pre-Existing Flaw
(2c)Pre-Existing Flawpλpλ
tλtλ
2
1
(2c)Pre-Existing Flaw
2
1
(2c)Pre-Existing Flaw
(2c)Pre-Existing Flaw
2
1
(2c)Pre-Existing Flaw
(2c)Pre-Existing Flaw
2
1
(2c)Pre-Existing Flaw
(2c)Pre-Existing Flawpλpλ
tλtλtλ
Figure 5. Wing-crack with plastic relaxation model proposed by Horii and Nemat-
Nasser29
for compressive failure.
normal stress and sliding friction resists this sliding motion. The critical shear
stress for wing-crack initiation and growth (assuming = 45crito and = 45
o):
1212critt1212
12*
Y
crit
11c112
1
λ (2)
Equation (2) can be used to predict the location and severity of compressive
damage in a solid by simply comparing it with the measured or estimated
maximum shear stress. Because failure is pressure dependent, the confining stress
must also be measured or estimated. For the problem under consideration in this
paper, it is assumed that the shear stress and confining stress generated during
dwell are given by the results of Hertz’s theory for frictionless contact between
axisymmetric linear-elastic bodies.
DWELL/PENETRATION TRANSITION MODEL DEVELOPMENT
Classic Hertzian Stress Distribution
The principal stresses and maximum shear stress in the elastic solid along the
center-line axis of contact (z-axis) arising due to the frictionless contact between
axisymmetric bodies with similar elastic moduli are given by31
:
2o
H1
az1
1
p (3a)
2
1
o
H3
o
H2
az12
1
z
atan
a
z11
pp (3b)
Ceramic Armor Materials by Design 563
0.5 1 1.5 2
0
0.5
1
1.5
2
2.5
z/a
Damage
* = 0.05
No Damage
No Damage
Hertz/
crit
* = 0.1
Figure 6. Plot of the center-line distribution and severity of damage for several
values of *.
2
1
o
H
az12
3
z
atan
a
z11
2
1
p (3c)
where po is the maximum interface normal stress, 2a is the contact diameter, is
Poisson’s ratio for the elastic solid, and superscript “H” signifies Hertz’s solution.
The general solution for stresses is reported by Lawn19
.
Predicted Spatial Distribution of Damage and Severity
Substitution of Equations 3(a) and 3(b) into Equation 2 yields an expression
for the critical shear stress as a function of normalized depth and material
parameters. Dividing Equation 3(c) by this expression yields Figure 6 where the
distribution of damage and its severity along the center-line axis is plotted for
several values of the ductility parameter* with po = Y, = 0.2, = 0, and
critt cλ = 0.1. Damage would only be expected whereH > crit. Thus, as can
be seen, damage would not be expected near the surface or at depths significantly
greater than the diameter of the contact. Near the surface, the stresses are high,
but the stress-state is more hydrostatic. It can also be seen from this figure, that
the expected severity of damage rises quite rapidly, reaching a maximum less than
one-half contact diameters below the surface. From this maximum, the severity
of damage gradually tapers off. It can also be noted that the distribution and
severity of damage is strongly effected by the ductility parameter. These
564 Ceramic Armor Materials by Design
predictions are entirely consistent with the localized damage shown in Figure 2,
as well as with previous observations18
.
Reconsidering the dwell/penetration transition velocity data shown in Figure
3(b) in light of these observations suggests the following possible explanations for
the upper and lower bound curves. Assuming the absence of shock-induced
damage, for ceramics such as boron carbide or Syndie, the comminuted region
forms and extends to the top surface. As a result, the damaged material within the
comminuted region becomes unconfined and is therefore easily displaced,
allowing penetration to occur. In the case of silicon carbide or titanium diboride,
this comminuted region forms, but does not extend to the top surface (i.e. it is
confined by the surrounding undamaged material). For penetration to occur, the
“undamaged” layer of ceramic separating the rod from the comminuted region
must fail. Therefore, for the silicon carbides and titanium diboride shown in
Figure 3(b), penetration is governed by the ductile failure of this relatively
undamaged layer. Consequently, the critical pressure for the dwell/penetration
transition would correspond approximately with 2.85 Y, that required to fully
indent a rigid, perfectly-plastic solid.
Critical Impact Pressure for the Dwell/Penetration Transition
Consideration of Equations 2 and 3 with the conditions that H = crit and z/a =
0, the critical mean pressure pm required to expand the damaged region to the top
surface is given by (for a Hertzian pressure distribution, the maximum po is equal
to 3/2 times the mean pressure pm31
):
critt
*
Y
m
c212321
93.0
85.2
p
λ (4)
Equation 4 relates the critical impact pressure pm to material properties and
characteristics. As such, if these material properties and characteristics are
known, Equation 4 can be used as a figure-of-merit and also used to predict the
dwell/penetration transition velocity. According to Lundberg et al.11
, the critical
impact pressure pm and the projectile dwell/penetration transition velocity Vp are
related by the following expression:
p
pYm
p
pp
K
27.3p211
K2V (5)
where Kp, p, and Yp are the bulk modulus, density, and yield strength of the
projectile, respectively. The predicted effect of the ductility parameter on the
dwell/penetration transition velocity is shown in Figure 7. The values assumed
Ceramic Armor Materials by Design 565
500
1000
1500
2000
2500
3000
5 10 15 20 25 3
Vp
(m/s)
Y (GPa)
* = 0.05
* = 0.1
* = 0.15
pm
= 2.85Y
0
Figure 7. Predicted effect of *
on the dwell/penetration transition velocity based
upon Equations 4 and 5.
for the ceramic were = 0.2, = 0, and critt cλ = 0.1. The values for the long-
rod projectile were p = 17.6 x 103 kg/m
3, Kp = 285 GPa, and Yp = 1.2 GPa
11.
According to this plot, a ceramic that possessed a dwell/penetration transition
velocity given by point A could be improved if its ductility parameter is
increased. This could be done by increasing the fracture toughness KIC,
decreasing the governing pre-existing flaw size 2c, decreasing the yield strength
(or hardness), or a combination of all three of these parameters.
Considering the Equation 4 as a figure-of-merit and in terms of the
dwell/penetration transition velocity, the following conditions would be possible:
If 185.p Ym 2 , the dwell/penetration transition velocity would
be given by the upper-bound curve in Figure 7.
If 185.2p Ym , the dwell/penetration transition velocity would
be less than that given by the upper-bound curve in Figure 7.
If a ceramic’s dwell/penetration transition velocity was below that predicted by
the upper-bound curve shown in Figure 7, the predicted change in this velocity
Vp for a change in mean impact pressure pm required to expand the damaged
region to the top surface is given by:
Am
m
p
2App
2App
Am
Ap
p
p
p
K2V1
Vp
V
V (6a)
566 Ceramic Armor Materials by Design
(6b) A
mi
mm ppp
criti
tiiii
*i
iY
im
c212321
93.0
85.2
p
λ (6c)
The maximum pm is given by:
(7) A
miYmaxm p85.2p
If the only difference between ceramic A and i is in their fracture toughness
values, then Equation 6(a) can be written as:
AIC
IC
p
2App
2App
Am
Ap
p
K
K
K2V1
Vp
V
V (8)
where KIC equals KiIC – K
AIC. Considering Equations 7 and 8, the maximum
change in fracture toughness is given by:
1p
85.2
K
K
Am
AY
AIC
maxIC (9)
For example, consider the boron carbide shown in Figure 3(b). The yield strength
is 15.8 GPa11
, while the mean impact pressure at 1450 m/s is approximately 23
GPa. Substitution into Equation 9 yields a maximum required change in fracture
toughness equal to 2.7 MPa*m1/2
assuming a base fracture toughness of 2.8
MPa*m1/2
(i.e. the KiIC = 5.5 MPa*m
1/2). The assumption that the only difference
between ceramics i and A are in their fracture toughness values was used in
deriving Equations 8 and 9. While it is acknowledged that actually producing a
ceramic with this quality would not necessarily be easy, the purpose was to
demonstrate the potential utility of the model for guiding ceramic developers.
SUMMARY
A hypothesis has been proposed to explain the upper bound and lower
bound(s) in the dwell/penetration transition velocity data of Lundberg et al.11
.
This hypothesis is based upon observations of the localized damage region in
ceramic tiles recovered from successful interface defeat experiments conducted
by Hauver et al.9. These observations led to the consideration of
micromechanical-descriptions for compressive failure of brittle solids. In
particular, the wing-crack model has been successfully used to explain the effect
Ceramic Armor Materials by Design 567
of hydrostatic stress on the failure mode of geological materials. Combining the
wing-crack model developed by Horii and Nemat-Nasser29
with the stress
distribution developed by Hertz19,31
for the normal, frictionless contact between
axisymmetric linear-elastic bodies led to the development of a model which
captured the essential features of both the distribution and severity of damage
within the localized damage region. It also provides a physically-based rationale
for the dwell/penetration transition data of Lundberg et al.11
.
The model indicates that for ceramic armor applications where part of the
defeat mechanism is dwell, the first consideration for the ceramic should be its
“hardness”. The second consideration should be its ductility parameter, or in
particular, its grain size and fracture toughness. Assuming constant physical and
mechanical properties, with the exception of fracture toughness, an equation was
derived that relates the change in fracture toughness to the change in
dwell/penetration transition velocity. Using boron carbide as an example, the
utility of the model for ceramic developers was shown.
While it was not discussed, the fracture toughness value that one should
consider is for cracks whose length-scale is less than or equal to the grain size.
For cracks of this size, the effect of residual stress will also be important32
.
Residual stress was not accounted for in this model. Finally, it must be mentioned
that this model is applicable to ceramics that do not have “soft” grain-boundary
phases. The shear strength of the “soft” grain-boundary phase could provide an
even lower critical shear stress than that given by Equation 2.
ACKNOWLEDGEMENTS
The author would like to thank Mr. William Bruchey, Mr. William Gooch,
Mr. George Hauver, Mr. Edward Horwath, Dr. Michael Normandia, Mr. Edward
Rapacki, Dr. Bryn James, and Dr. Patrick Lundberg for sharing their knowledge
on the phenomenon of dwell and interface defeat. The author would also like to
thank Professor Marc Meyers for sharing his knowledge and insight into the
physical mechanisms that govern the dynamic behavior of ceramics.
REFERENCES1W.E. Snowden, “High Performance Ceramics for Armor Applications: A
Historical Perspective,” in Proceedings of the Symposium on Ceramic Armor
Materials by Design, PAC RIM 4, Wailea, Maui, HI, November 4-8, 2001.2B. Matchen, “Applications of Ceramics in Armor Products,” Key Eng. Mat.,
122-124 333-42 (1996). 3Z. Rozenberg and Y. Yeshurun, “The Relation Between Ballistic Efficiency
and Compressive Strength of Ceramic Tiles,” Int. J. Impact Eng., 7 [3] 357-62
(1988).
568 Ceramic Armor Materials by Design
4J. Sternberg, “Material Properties Determining the Resistance of Ceramics to
High Velocity Penetration,” J. Appl. Phys., 65 [9] 3417-424 (1989). 5M.A. Adams, “Theory and Experimental Test Methods for Evaluating
Ceramic Armor Components,” in Proceedings of the Symposium on Ceramic
Armor Materials by Design, PAC RIM 4, Wailea, Maui, HI, November 4-8, 2001. 6M.J. Normandia and W.A. Gooch, “An Overview of Ballistic Testing
Methods of Ceramic Materials,” in Proceedings of the Symposium on Ceramic
Armor Materials by Design, PAC RIM 4, Wailea, Maui, HI, November 4-8, 2001. 7G.E. Hauver and A. Melani, “Behavior During Penetration of Long Rods
(U)”; pp. 149-160 in Proceedings of the Second BRL Topical Symposium:
Experimental Research and Modeling Support, Ballistic Research Laboratory,
Aberdeen Proving Ground, MD, May 24, 1988.8G.E. Hauver, P.H. Netherwood, R.F. Benck, and L.J. Kecskes, “Ballistic
Performance of Ceramic Targets”; pp. 23-34 in Proceedings of the 13th
Army
Symposium on Solid Mechanics, Plymouth, MA, August 17-19, 1993. 9G.E. Hauver, P.H. Netherwood, R.F. Benck, and E.J. Rapacki, “Interface
Defeat of Long-Rod Projectiles by Ceramic Armor,” ARL Technical Report, in
progress.10
P. Lundberg, L. Holmgren, and B. Janzon, Ballistics ’98, in Proceedings of
the Seventeenth International Symposium on Ballistics, Midrand, South Africa,
March 1998, 3 251- (1998).11
P. Lundberg, R. Renstrom, and B. Lundberg, “Impact of Metallic Projectiles
on Ceramic Targets: Transition Between Interface Defeat and Penetration,” Int. J.
Impact Eng., 24 259-75 (2000).12
J.E. Field, “High-Speed Photography,” Contemp. Phys., 24 [5] 439-59
(1983).13
D.A. Shockey, A.H. Marchand, S.R. Skagg, G.E. Cort, M.W. Burkett, and
R. Parker, “Failure Phenomenology of Confined Ceramic Targets and Impacting
Rods,” Int. J. Impact Eng., 9 [3] 263-75 (1990). 14
Y. Tanabe, T. Saitoh, O. Wada, H. Tamura, and A.B. Sawaoka, “An
Overview of Impact Damages in Ceramic Materials – For Impact Velocity Below
2 km/s,” Report of the Research Laboratory of Engineering Materials, Tokyo
Institute of Technology, 19, 1994. 15
E. Strassburger, H. Senf, C. Denoual, P. Riou, and C. Cottenot, J. Phys. IV
France, 7 [c3] 909-14 (1997). 16
D. Sherman, “Impact Failure Mechanisms in Alumina Tiles on Finite
Thickness Support and the Effect of Confinement,” Int. J. Impact Eng., 24 312-28
(2000).17
C.J. Shih, Dynamic Deformation of Silicon Carbide, Ph’d Dissertation,
UMI, Dissertation Information Service, 1998, 331 pp.
Ceramic Armor Materials by Design 569
18J.C. LaSalvia, E.J. Horwath, E.J. Rapacki, C.J. Shih, and M.A. Meyers,
“Microstructural and Micromechanical Aspects of Ceramic/Long-Rod Projectile
Interactions: Dwell/Penetration Transitions”; pp.437-46 in Fundamental Issues
and Applications of Shock-Wave and High-Strain-Rate Phenomena, ed. K.P.
Staudhammer, L.E. Murr, and M.A. Meyers, Elsevier Science, New York, 2001. 19
B.R. Lawn, “Indentation of Ceramics with Spheres: A Century after Hertz,”
J. Am. Ceram. Soc., 81 [8] 1977-94 (1998). 20
D.A. Shockey, D.J. Rowcliffe, K.C. Dao, and L. Seaman, “Particle Imapct
Damage in Silicon Nitride,” J. Am. Ceram. Soc., 73 [6] 1613-19 (1990). 21
D.K. Kim, C-S. Lee, C.W. Kim, and S.N. Chang, “Indentation Damage
Behavior of Armor Ceramics,” in Proceedings of the Symposium on Ceramic
Armor Materials by Design, PAC RIM 4, Wailea, Maui, HI, November 4-8, 2001. 22
K. Zeng, E. Soderlund, A.E. Giannakopoulos, and D.J. Rowcliffe,
“Controlled Indentation: A General Approach to Determine Mechanical
Properties of Brittle Materials,” Acta Mat., 44 [3] 1127-41 (1996). 23
J. Alcala, A.E. Gainnakopoulos, and S. Suresh, “Continuous Measurements
of Load-Penetration Curves with Spherical Microindenters and the Estimation of
Mechanical Properties,” J. Mat. Res., 13 [5] 1390-1400 (1998). 24
Yu. V. Milman and S.I. Chugunova, “Mechanical Properties, Indentation
and Dynamic Yield Stress of Ceramic Targets,” Int. J. Impact Eng., 23 629-38
(1999).25
Fracture in Compression of Brittle Solids, National Materials Advisory
Board, Report 404, National Academy Press, August 1983, 70 pp. 26
B. Cotterell, “Brittle Fracture in Compression,” Int. J. Fracture, 8 [2] 195-
208 (1972). 27
E.Z. Lajtai, “Brittle Fracture in Compression,” Int. J. Fracture, 10 [4] 525-
36 (1974). 28
M.F. Ashby and S.D. Hallam, “The Failure of Brittle Solids Containing
Small Cracks Under Compressive Stress States,” Acta Mat., 34 [3] 497-510
(1986).29
H. Horii and S. Nemat-Nasser, “Brittle Failure in Compression: Splitting,
Faulting, and Brittle-Ductile Transition,” Phil. Trans. R. Soc. London A, 319 337-
374 (1986). 30
M.F. Ashby and C.G. Sammis, “The Damage Mechanics of Brittle Solids in
Compression,” PAGEOPH, 133 [3] 489-521 (1990). 31
K.L. Johnson, Contact Mechanics, Cambridge University Press, 1985,
452pp.32
S.J. Bennison and B.R. Lawn, “Role of Interfacial Grain-Bridging Sliding
Friction in the Crack Resistance and Strength Properties of Nontransforming
Ceramics,” Acta Mat., 37 7659-71 (1989).
570 Ceramic Armor Materials by Design
Transparent Armor
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TRANSPARENT ARMOR MATERIALS: NEEDS AND REQUIREMENTS
Parimal J. Patel and Gary A. Gilde
Army Research Laboratory,
Weapons and Materials Research Directorate
Attn: AMSRL-WM-MC
Aberdeen Proving Ground, MD 21005
ABSTRACT
There has been interest in improving transparent armor for use in Army
vehicles. Future combat and non-combat environments will require lightweight,
threat adjustable, multifunctional, and affordable armor. Significant
improvements can be achieved through insertion of new materials. However, an
emphasis must be placed on user needs and requirements in addition to
improvements in ballistic performance. Current glass/polycarbonate technologies
are not expected to meet the increased requirements for transparent protection.
Results over the past few years indicate that the use of transparent crystalline
ceramics and advanced polymers greatly improve the performance of a system.
An overview of user requirements, applications, and current efforts in transparent
armor will be discussed.
Keywords: transparent, ceramic, armor, aluminum oxynitride, spinel,
sapphire, polycarbonate, polyurethane
INTRODUCTION
Transparent Armor Requirements
Transparent armor is a system constructed of different materials that are
designed to defeat a particular threat or range of threats. The threats targeted are
dependent on the envisioned combat or non-combat scenarios. There are also
threat requirements for "operations other than war" where ballistic protection is
required. Though a system is designed for a particular threat, there are general
requirements common to most transparent armor systems. The paramount
requirement for a transparent armor system is the defeat of a designated threat.
The system must also provide a multi-hit capability with minimized distortion of
Ceramic Armor Materials by Design 573
To the extent authorized under the laws of the United States of America, all copyright interests in this publication are the propertyof The American Ceramic Society. Any duplication, reproduction, or republication of this publication or any part thereof, withoutthe express written consent of The American Ceramic Society or fee paid to the Copyright Clearance Center, is prohibited.
surrounding areas of the first hit. They must be transparent in the wavelengths of
interest, ranging in the UV, visible, and infrared frequencies. Other requirements
for transparent armor windows are that they are night vision compatible, and they
are affordable based on cost-performance models.1
For future land and air platforms, weight is a critical parameter that must be
minimized. Space efficiency can also be quite important for certain applications.
The system must be large enough to be useful. For example, a 6 inch square
armor plate may be useful as a face shield but would not be very useful as a truck
window. The size must be large enough for the user to perform their duties
appropriately. Baseline transparent armor systems generally rely on plastics,
plastic-plastic laminates, and glass/plastic laminates. The systems work and offer
protection for the threats they are designed for. As the defined threat become
more lethal, these systems no longer perform adequately. A simple solution that
increases the ballistic performance of a window is increasing the thickness of the
window. The material and design costs are thus, increased incrementally. For
many applications, very thick armor systems are not practical solutions, even if
they defeat the threat. Thick windows may be impractical for a few reasons. One
reason is due to the increased weight associated with thicker materials. Another
reason is the space limitations in many vehicles. Finally, thick sections of
transparent armor have greater optical distortion than thinner sections, reducing
the transparency. Therefore, new materials that are thinner, lightweight, and offer
better ballistic performance are sought. Affordability is a critical metric for
evaluating all armor systems and can be the limiting factor for given applications.
There are many methods to measure the ballistic performance of a system.
Several experimental techniques have been developed to aid in comparative
studies of armor systems. One of these tests, a V50 test2, was used to measure
ballistic performance for the systems mentioned in this paper.
A transparent armor system is comprised of many layers, joined by polymeric
interlayers. The front face is usually a hard face material that is designed to break
up or deform the projectile upon impact. The sequential plys are added to provide
additional resistance to penetration. These materials can be the same as or
different than the front ply material. An interlayer to join the two plates separates
the plys and provides a transition between two materials that may have thermal
expansion mismatches such as a glass and a polymer. The purpose of this
interlayer is to mitigate the stresses from thermal expansion mismatches, as well
as to stop crack propagation from ceramic to polymer during processing.
The armor system can be engineered to provide different levels of protection.
In addition to defeating the threat with multi-hit capability, the mass and space
efficiency should be optimized for a given application. The variables that can be
changed are plate material, thickness of plys, interlayer hardness, interlayer
thickness, number of plys and the order of constituent materials.
574 Ceramic Armor Materials by Design
MATERIALS USED FOR TRANSPARENT ARMOR
Polymeric Materials
The most common plastic used for transparent armor applications is
polycarbonate. Polycarbonate offers excellent ballistic protection against small
fragments. Polycarbonate is an inexpensive material that is easily formed or
molded. Polycarbonate is used in applications such as the sun, wind, and dust
(SWD) goggles, spectacles, visors, face shields and laser protection goggles.
Polycarbonate is also used as a backing material for advanced threats.
Polycarbonate is more effective in the thin dimensions required for individual
protection than in the thicker sections required for vehicle protection. Though the
material is adequate for many applications, the search for lighter weight materials
has led to investigations into other polymeric materials such as transparent nylons,
polyurethane, and acrylics.3,4
The limiting factor for use of other transparent
polymeric materials is their durability and their optical properties. Improvement
in these properties would warrant an investigation into the ballistic properties of
the material.
There have been efforts to improve the properties of polyurethane. Simula
Technologies Inc. has recently introduced a new family of polyurethanes with
improved optical properties. These materials are marketed and sold by Simula
Polymer Systems Inc. Sim 2003 and Sim 1802 are both thermoset plastics that are
produced by casting or liquid injection molding. Sim 1802 is harder and more
brittle than Sim 2003. Due to their physical properties, Sim 2003 is a viable
candidate to replace polycarbonate as a riot visor or as a backing material. Sim
1802 is a better candidate for front or hard-face material. These improvements in
polyurethane have led to an investigation into these materials for face-shield
applications as will be discussed in the “Applications” section.
Glasses and Glass-Ceramics
There are several glasses that are utilized in transparent armor. Normal plate
glass (soda-lime-silica) is the most common glass used due to its low cost, but
greater requirements for optical properties and ballistic performance generate the
need for new materials. There are many different glasses including borosilicate
glasses and fused silica that can be used. Glasses can be strengthened using
chemical or thermal treatments. Controlled crystallization of certain glass systems
can also produce transparent glass-ceramics. TransArm, a lithium disilicate based
glass-ceramic is produced by Alstom++
for use in transparent armor systems.5
Simula Technologies, 10016 South 51st Street, Phoenix, AZ, 85044 ++
Alstom UK Ltd., Research & Technology Centre Stafford, Staffordshire, ST174LN, England.
Ceramic Armor Materials by Design 575
Glasses and glass-ceramics have the overall advantage of having lower cost than
most other ceramics materials, and the ability to be produced in curved shapes and
large sheets.
Transparent Crystalline Ceramics
For advanced threats, transparent crystalline ceramics are used to defeat the
projectiles. However, there are not many candidate ceramic materials, however,
that are transparent. The three major candidates are aluminum oxynitride (AlON),
magnesium aluminate spinel (spinel), and single crystal aluminum oxide
(sapphire). There are advantages and disadvantages to each material.
Aluminum Oxynitride Spinel (Al23O27N5): One of the leading candidates for
transparent armor is aluminum oxynitride (AlON). It is produced by Raytheon
Corporation+
and marketed under the trade name Raytran. The incorporation of
nitrogen into an aluminum oxide stabilizes a spinel phase. Due to its cubic crystal
structure, AlON is an isotropic material that can be produced transparent as a
polycrystalline material. A polycrystalline material can also be produced in
complex geometries using conventional ceramic forming techniques such as
pressing and slip casting. The green body is processed to transparency and
polished. Some properties of AlON are listed on Table I. The limitations of AlON
are its high cost and the sizes that are currently available. Raytheon is currently
investigating the scale-up and costreduction of aluminum oxynitride.
Raytheon Corp has produced an 11in. x 11in. curved AlON window (Figure
1A). The Air Force Research Laboratory (AFRL) is currently funding Raytheon
to investigate cost reduction of AlON to produce larger windows. This will allow
Raytheon to scale-up AlON so that it can be produced in large sizes at reasonable
costs. Additionally, funding is sought to address the equipment issues to produce
very large size plates.
Concurrently, the U. S. Army Research Laboratory (USARL) is investigating
transient liquid phase sintering of aluminum oxynitride to reduce processing
costs.6
A reaction sintering technique with the aid of a reactive liquid is the focus
of the research. Small samples (Figure 1B) with a transmission of 85% and a haze
of 14% have been produced. The reduction of the haze and size scale-up are the
immediate objectives of the program. ARL also has a Small Business Innovative
Research (SBIR) program for processes that can produce affordable aluminum
oxynitride powders using scalable methods to reduce the cost of the raw
materials.
+ Raytheon Electronic Systems, Lexington Laboratory, 131 Spring Street, Lexington, MA 02421
576 Ceramic Armor Materials by Design
Figure 1: Photographs of aluminum oxynitride produced by Raytheon and ARL
Magnesium Aluminate Spinel (MgAl2O4): Spinel is a transparent ceramic that
has a cubic crystal structure and can be transparent in its polycrystalline form.
Spinel produced by sinter/HIP, hot pressing, and hot-press/Hot Isostatic Pressing
(HIP) has yielded transparent materials.7,8
The use of a hot isostatic press has
been shown to improve the optical and physical properties of spinel.8
Table I
shows the properties of spinel. Spinel offers some processing advantages over
AlON. Spinel powder is available from commercial powder manufacturers while
AlON powders are proprietary to Raytheon. Spinel is also processed at much
lower temperatures that AlON. The optical properties are better than AlON, with
its IR cut-off at 6 microns compared to 5.5 microns for AlON and 6 microns for
sapphire, respectively.9
Though spinel shows promise for many applications, it is
not available in bulk form from any manufacturer, but there are efforts to
commercialize spinel.
Table I: Selected mechanical properties of AlON and spinel
AlON Spinel
Density g/cm3
3.67 3.58
Elastic Modulus GPa 315 277
Mean Flexure Strength MPa 228 241
Weibull Modulus 8.7 19.5
Fracture Toughness MPa m 2.4±0.11 1.72±0.06
Knoop Hardness (HK2) GPa 13.8±0.3 12.1±0.2
Ceramic Composite Inc.+
is currently investigating hot pressing of magnesium
aluminate spinel under a Phase II SBIR sponsored by the Army Research
Laboratory. Previous investigations have studied sinter-hot isostatic pressing
(HIP) techniques. Hot pressing was chosen for this program as the processing
Ceramic Armor Materials by Design 577
technique based on comparative analysis of the several processing techniques for
producing spinel.8
The research has focused on hot pressing with an additive and
hot-press/hot isostatic pressing (HIP). Hot pressing has been shown to be a
successful technique to produce transparent parts. Figure 2 is a four-inch
diameter, 0.44-inch thick spinel plate that has been produced using this technique.
The plate has an 83 percent transmission with 9.32 percent haze. Scale-up to
twenty inch parts is underway using the hot-press technique. Subsequent HIPing
has been shown to improve the optical properties and mechanical properties of
spinel.7
Hipping is generally not cost-effective and its use should be minimized.
However, the improvement in the mechanical and optical properties may deem
HIPing necessary for given applications.
Figure 2: A hot pressed 4-inch diameter, 0.44" thick spinel plate produced at ARL
Single Crystal Aluminum Oxide (Sapphire - Al2O3): Polycrystalline aluminum
oxide is an armor ceramic material that is used in opaque armor systems.
Aluminum oxide is transparent when produced in single crystal form. The
material is grown using single crystal growth techniques such as HEM10
by
Crystal Systems Inc.+
or edge-defined film-fed growth (EFG)11
by Saphikon.++
The crystal structure of sapphire is rhombohedral and its properties are
anisotropic and vary with crystallographic orientation. Sapphire is currently the
most mature transparent ceramic and is available from several manufacturers. The
cost is high due to the processing temperature involved and machining costs to cut
parts out of single crystal boules. Sapphire is a very high strength material, but the
strength is very dependent on the surface finish.12
There are current programs to
scale-up sapphire grown by the HEM and EFG processes.
+ Crystal Systems, Inc., 27 Congress St., Salem, MA 01970
++ Saphikon, 33 Powers St., Milford, NH 03055
578 Ceramic Armor Materials by Design
Another manufacturer of sapphire is Saphikon, Inc., which produces
transparent sapphire using an edge, defined growth technique. The process size
limitation is currently at 0.25 in. thick, in 10 in. x 10 in. sheets. The Army
Research Laboratory is currently investigating use of this material for transparent
armor systems using synergistic approaches in laminate design and construction.
The current objective is to determine a baseline of glass/plastic and
ceramic/plastic against the specified threat. Once the baseline is completed,
sapphire will be tested in different constructions and compared to the baseline.
Manufacturing: Scale-up to larger size poses several problems. The large
sizes generally cost more to produce due to the difficulty in scale-up. Also, larges
plates are more difficult to polish than smaller plates. Materials Systems Inc.+
is
investigating bonding sapphire plates using proprietary glass and glass-ceramic
bonding materials.13
To date, bonds have been produced that are 70 percent of the
strength of unbonded material.13
This innovative technique offers the ability to
make very large windows that may not be achievable in monolithic parts due to
lack of capital equipment. A 12.4" x 18.9" window bonded is shown in Figure 3.
The limitation of this process is the presence of bond lines that are presently
visible. There are efforts to remove or minimize these visual effects.
Figure 3: Ground sapphire bonded together by MSI to form a 12.4" by 18.9"
window
Machining and Polishing: Regardless of the ceramic material utilized,
machining and polishing costs are significant. The high hardness of AlON, spinel,
and sapphire require diamond grinding and polishing media. The finishing
+ Materials Systems Inc., 521 Great Road, Littleton, MA, 01460
Ceramic Armor Materials by Design 579
process times are also quite long. Finishing costs can be as much as 50 percent of
the final cost of the materials. These costs are greater for curved windows.
There are some programs to reduce the costs of machining and polishing. The
Center for Optics Manufacturing+
is investigating advanced grinding and
polishing techniques for optics. Their processing has been shown to remove
AlON, sapphire, and SiC at removal rates of 3 um/min, 1.5 um/min, and
0.5um/min, respectively.14
Figure 4: Low cost alternative to polishing developed by MSI. The left
photograph is with no window while the photo on the right is the view looking
through coated ground sapphire plate.
The Army Research Laboratory is also is looking for low cost solutions to
polishing. The USARL has recently sponsored a Phase II SBIR to address low
cost alternatives for polishing. Materials Science Inc. of Littleton, MA, is
investigating various treatments on ground sapphire to make it transparent without
polishing. The initial results have been very successful, as can be seen in Figure 4.
The technique eliminates the final polishing step thus saving significant amounts
of time and cost in producing the transparent ceramic.
APPLICATIONS AND REQUIREMENTS
Common military applications for transparent armor are ground vehicle
protection, air vehicle protection, personnel protection, and equipment (sensor)
protection. There are also commercial applications such as riot gear, face shield,
security glass, armored cars and armored vehicles.
Personnel Protection
There are several applications of advanced transparent armor systems for
personnel protection. Personnel protection utilizes transparent armor against small
arms threats and fragments, such as high velocity, rocks and bottles. Goggles are
+ Center for Optics Manufacturing, 240 East River Road, Rochester, New York, 14623
580 Ceramic Armor Materials by Design
required for protections against the sun, wind and dust and in some cases, lasers.
Increased use of military forces for "operations other than war" highlights the
need to protect forces involved in these peacekeeping missions. For these
operations, protective equipment such as riot gear is needed. Laser threats are also
significant, and protective materials and coatings are sought for these
applications. Once again, improved ballistic protection and lighter weight are the
major objectives and cost is a significant factor.
Face shields: Personnel protection for facial protection is one Army
application that requires transparent armor. The Army Research Laboratory has
completed a program to improve the current visor design.15
The two end items
identified for improvement were the riot visor and an explosive ordnance (EOD)
visor. The goal for the riot visor was to improve the ballistic performance by 30
percent without increasing the weight of the system. The overall goal for the EOD
visor was to reduce the weight of the visor by 30 percent while providing equal
protection.
Riot Visor: The riot visor is made from injection-molded polycarbonate that
has an areal density of 1.55 lb/ft2. The visor is designed to protect against large,
low-velocity projectiles such as rocks and bottles, as well as, from small, high
velocity fragments. Since the goal of this program was to improve the ballistic
performance without increasing the weight, an all-polymer solution was sought.
Previous investigations16,17
in the 1970's had shown the promise of polyurethane
as an armor material, but the optical properties were not sufficient for a
transparent armor material. Improvements in the optical properties of the
polyurethane by Simula warranted a ballistic evaluation.
Ballistic testing was conducted for the riot visor against a 0.22 cal fragment
simulating projectile (FSP).1
A helium gas gun was used for velocities below 2000
ft/sec and a 22 inch-long, 0.223 barrel with a 1:12 twist was used for velocities
above 2000 ft/sec. The results of the testing showed that the polyurethane (SIM
2003) behaves better than either polycarbonate or acrylic (PMMA). Overall, the
polyurethane performed 30-35 percent better than polycarbonate on an equal
weight basis. The conclusion was that with the improved optical properties of the
SIM 2003, this material would be an excellent replacement for polycarbonate to
reduce the weight of the system.
Explosive Ordnance Visor (EOD:) The other objective for the ARL program was
to reduce the weight of EOD visors. The goal is to reduce the areal density of the
current system using different materials and constructions. Several constructions
were investigated, including plastic/plastic laminates, glass/plastic laminates, and
Ceramic Armor Materials by Design 581
glass-ceramic/plastic laminates.15
The plastic hard-face did not deform the FSP,
while the glass and glass-ceramics were able to deform the FSP. Many of the
constructions were better in weight than the current system weight of 4.27 lb/ft2.
The use of a polyurethane (Sim 2003) increased the performance of the system.
The optimum constructions used fused silica, Vycor, or TransArm, a transparent
glass-ceramic produced by GEC Alsthom.
The ballistic data obtained in this investigation can be used for comparative
purposes in designing a visor for use against the FSP threat for the range of areal
densities tested. Other considerations are cost, availability, and manufacturability,
for which there are trade-offs. For example, in visor applications, TransArm,
Vycor, and fused silica performed well. TransArm is currently more expensive
than fused silica. However, TransArm can be easily produced in curved shapes.
Currently, it is difficult to obtain fused silica in a curved shape of a visor. Thus,
while fused silica would be a lower cost solution that performs better (optically
and ballistically) it may not be used for visor applications until the manufacturing
problem of producing fused silica in curved shapes is overcome.
Ground Vehicles
Ground vehicle protection is required for equipment that is used on the
battlefield, such as HMMWVs, tanks, trucks, and resupply vehicles. Transparent
armor is necessary for the windshield and side windows. There are several general
requirements for these applications.18
One critical requirement is the ability to
withstand multiple hits since most threat weapons are automatic or
semiautomatic. The windows must also be full size so that the vehicle can be
operated in the manner in which it was designed. A small window on a truck can
increase ballistic survivability but can reduce operational safety if the driver does
not have an appropriate field of view. The windows also need to be durable and
withstand normal wear in non-combat situations and from user damage.
The fielded systems fulfill these requirements with varying degrees of
success. There are some requirements that future transparent armor systems need
to address.18
There is an overall requirement for future Army systems to be
lighter. The weight of a transparent armor system is a parasitic weight for a
vehicle. The added weight of a transparent armor appliqué can be significant,
often requiring a beefed up suspension and drive train to maintain the vehicles
performance capability. These upgrades also add weight to the system. Any
weight savings improves the ability to bring the vehicle into theater. Reduction in
weight increases the payload capacity for tactical vehicles and thus increases
operational capabilities. Thinner armor systems are also required for similar
reasons. Thinner windows can increase the cabin volume. Future systems also
need to be compatible with night vision goggle equipment while offering laser
protection.
582 Ceramic Armor Materials by Design
Due to their size and shape, windows are constructed of glass and plastic. The
major drive for new windows for these applications is lower weight and improved
ballistic protection. Due to the number of these vehicles in service, the sizes of the
windshields needed and the costs, improved glasses, glass ceramics and polymers
are the materials of choice for these applications. Glass compositional variations,
chemical strengthening, or controlled crystallization can improve the ballistic
properties. Glasses can also be produced in large sizes and curved shapes. Most
importantly, glasses can be produced to provide incremental ballistic performance
and incremental cost.
For advanced threats, the weight of glass/plastic becomes prohibitively heavy
and thick. The use of a transparent ceramic as a front-ply has been shown to
improve the ballistic performance and reduce the weight of the system. The use of
a ceramic front ply can reduce the areal density by as much as 65 percent. This is
a significant weight savings over the state-of-the art. The ballistic performance of
these transparent ceramics offers great potential for weight savings on future
vehicles. Currently there are some challenges that must be overcome for these
materials to be utilized. The major limitations are cost, sizes available, and
curvature of the plates. There are several programs addressing these issues at
USARL and elsewhere as was described in the "Materials" section.
Aircraft
Helicopters and other aircraft used in combat or in support roles require
transparent armor. Applications include windshields, blast shields, lookdown
windows, and sensor protection. The general requirements for these systems are
similar to those for ground vehicles, though the importance of the requirements
varies. Weight is a critical factor for these applications. The current transparent
armor weight is the limiting factor for increasing ballistic protection. Heavier
vehicles use more fuel, are more difficult to move into theater, and reduce
maneuverability. The shields for aircraft applications need to be full size and
curved.
Electromagnetic windows
Many of the ceramic materials that are of interest for transparent armor
solutions are also applicable to electromagnetic (EM) windows. However, there
are many EM window applications where visible transparency is not critical. EM
window applications include radomes, IR domes, sensor protection, laser
windows,19
and multi-spectral windows. The requirements for these windows vary
greatly. There are some required properties mutual to many of the applications.
The optical properties are extremely important for window applications. The
transmission window and related cut-offs (uv, IR) control the electromagnetic
regime where the window is operational. Other properties of interest are abrasion
Ceramic Armor Materials by Design 583
resistance, strength, and the thermal properties. The thermal stability of the
materials properties are also critical if the material will be heated as in the case of
missile windows during flight.
Commercial Applications
Many of these systems utilized for military applications would also have use
in commercial systems such as law enforcement protection visors, riot gear, and
windows in commercial car, trucks, and busses, as well as architectural
requirements in certain buildings. The desire for armored automobiles for
personal use is also growing.
CONCLUSIONS
There is a general push to reduce the weight of military systems. Increased
weight reduces maneuverability, transportability, and increases operation costs.
One approach to reduce weight is to reduce the weight of armor systems. In
addition to reduction of weight, new systems are required to defeat more
advanced threats and to perform in combat and non-combat scenarios.
The history of transparent armor has shown significant advances as new
materials are introduced into the marketplace. The current thrust into lighter
systems is also based on advances in materials technology. Advances in
polymeric materials utilized for transparent armor systems have led to a renewed
interest in these materials to reduce the overall weight of armor systems.
Polyurethane has been shown to improve the performance as compared to
polycarbonate backing. Transparent ceramics have been shown to offer significant
ballistic protection with reduced weights over conventional glass/plastic systems.
Advances in the processes of these ceramics and scale-up have lead to increased
interest in using these materials for transparent armor applications.
There are several programs that are investigating the cost reduction and scale-
up of these materials. Successful outcomes from these programs should initiate
their use for armor applications and fulfill the requirements to reduce weight on
Army systems.
REFERENCES
1. P.J. Patel.; G.A.Gilde.; P.G. Dehmer, J.W. McCauley; "Transparent ceramics
for armor and EM window applications," PROC. SPIE Vol. 4102, p1-14,
Inorganic Optical Materials II, Alexander J. Marker; Eugene G. Arthurs; Eds.,
10/2000.
2. U.S. Department of Defense, "V50 Ballistic Test for Armor", MIL-STD662,
18 December 1997.
584 Ceramic Armor Materials by Design
3. A. L Lastnik,., M B Cleavly, J. B Brown,., "Development and fabrication of
polycarbonate eyeshield for the U.S. Army's Flyer's Helmet, TR-71-3-CE,
U.S. Army Natick Laboratories, Natick, MA, June 1970.
4. F. P Meyer, R Sacher, "Solarization effects on the materials employed in the
ballistic/laser eye protection spectacle system (B/LEPS), Interim Letter
Report, U.S. Army Materials Technology Laboratory, Watertown, MA, May,
1991.
5. A. R Hyde, J. G Darrant, "TRANSARM-Improved transparent armour,"
Proceedings of DARPA/ARL/ARO Transparent Armor Materials Workshop,
November 16-17, 1998, Annapolis, MD.
6. P.J Patel, G. A Gilde, J. W McCauley,., "Transient liquid phase sintering of
aluminum oxynitride (AlON), Army Research Laboratory Patent Disclosure
6-00, May 2000.
7. D.W Roy, J. L Hastert,L. E Coubrough, K. E Green, A Trujillo, "Method for
producing transparent polycrystalline body with high ultraviolet
transmittance," U.S. Patent # 5244849, September 14, 1993.
8. G.A. Gilde,P.J. Patel, M.Patterson, "A comparison of hot-pressing, rate
controlled sintering, and microwave sintering of magnesium aluminate spinel
for optical applications," Proceedings of SPIE Conference on Window and
Dome Technologies and Materials VI, Randal Tustison, SPIE Vol.3705, 94-
104, SPIE, Washington, April 1999.
9. D.C. Harris, Infrared window and dome materials, SPIE, Washington, pg.
32,1992.
10. Schmid, F., Viechnicki, D., J., Growth of Sapphire Disks from the Melt by a
Gradient Furnace Technique, J. Am. Ceram. Soc., 53, 528-29 1970.
11. H.E. Labelle, EFG, The Invention and Application to Sapphire Growth," J.
Cryst. Growth, 50, 8-17, 1980.
12. P.J. Patel, J.J. Swab,G.A. Gilde, Fracture properties and behavior of
transparent ceramics, PROC. SPIE Vol. 4102, p1-14, Inorganic Optical
Materials II, Alexander J. Marker; Eugene G. Arthurs; Eds., 10/2000.
13. P. McGuire, R. Gentilman, B. Pazol, J, Askinazi, J. Locher, "Mulitpane large
area and doubly-curved sapphire windows," Proceedings of the 8th DOD
Electromagnetic Windows Symposium, 27 April 2000.
14. H Policove,., "State of the Art in optical finishing," Proceedings of
DARPA/ARL/ARO Transparent Armor Materials Workshop, November 16-
17, 1998, Annapolis, MD.
15. P.J. Dehmer,., M. Klusewitz, "Proceedings of 8th DoD Electromagnetic
Windows Symposium at the USAF Academy, 24-27 April 2000"
16. R.W Lewis, and G.R Parsons, Ballistic Performance of Transparent Materials
for Eye Protection, AMMRC-TR-72-36, U.S. Army Material and Mechanics
Research Center, Watertown, MA, November, 1972.
Ceramic Armor Materials by Design 585
17. M.E Roylance,., and Lewis, R.W., Development of Transparent polymers for
Armor, AMMRC-TR-72-23, U.S. Army Material and Mechanics Research
Center, Watertown, MA, July, 1972.
18. R Gonzalez, G.J Wolfe, Ballistic Transparencies for Ground Vehicles,
Proceedings of DARPA/ARL/ARO Transparent Armor Materials Workshop,
November 16-17, 1998, Annapolis, MD.
19. R. A Beyer, H Kerwien, "Evaluation of AlON for cannon window
application," Proceedings of SPIE Conference on Window and Dome
Technologies and Materials VI, Randal Tustison, SPIE Vol.3705, 113-118,
SPIE, Washington, April 1999.
586 Ceramic Armor Materials by Design
MICROWAVE REACTIVE SINTERING TO FULLY TRANSPARENT
ALUMINUM OXYNITRIDE (ALON) CERAMICS
Dinesh Agrawal, Jiping Cheng, and Rustum Roy
Materials Research Institute
The Pennsylvania State University
University Park, PA 16802, USA
ABSTRACT
Fully transparent aluminum oxynitride (ALON) ceramic has been developed
by a single-step microwave sintering method. Starting with -alumina and
aluminum nitride powder mixture, the compacted pellets were microwave sintered
under an ambient pressure of pure nitrogen. It was found that single ALON phase
formed at 1650 C in 60 minutes by microwave processing, and the fully dense
and highly transparent ALON samples were made at 1800 C with residence time
of 60 minutes.
INTRDUCTION
Aluminum oxynitride (ALON) has an approximate composition of
Al23O27N5 (9Al2O3 5AlN). ALON can be sintered to fully transparent ceramic
material having mechanical and optical properties similar to those of sapphire
with the advantages of an isotropic cubic crystal structure. The transmission
range of ALON can extend from 0.2 m in the UV through the visible to 6.0 m
in the infrared, which makes it a very useful material for many electromagnetic
window applications. Combined with the high strength and high hardness, ALON
is an ideal material for transparent armor product [1,2].
The conventional fabrication of transparent ALON ceramics involves using
pre-synthesized ALON powder to form a green body, followed by sintering in a
nitrogen atmosphere at high temperatures (>1850 C) for extended period (20-100
hours) and often requires hot pressing [3]. A single-step preparation method was
also tried to make transparent ALON ceramics, by mixing Al2O3 and AlN
powders and subsequent reactive sintering at 1850 C for 1 hour at 3 bar nitrogen
atmosphere. But the sintered body in this case was translucent [4,5].
Microwave sintering is a novel sintering process that is fundamentally
different from the conventional sintering process. In conventional sintering, the
Ceramic Armor Materials by Design 587
To the extent authorized under the laws of the United States of America, all copyright interests in this publication are the propertyof The American Ceramic Society. Any duplication, reproduction, or republication of this publication or any part thereof, withoutthe express written consent of The American Ceramic Society or fee paid to the Copyright Clearance Center, is prohibited.
sintering driven force, temperature, is generated by external heating elements (in
resistance heating) and then is transferred to the samples via radiation, conduction
and convection. In microwave process, the processing materials themselves
absorb microwave power and then convert microwave energy in to heat within the
sample volume itself, and hence the heating is very rapid and uniform. The
microwave processing of materials has major advantages of higher energy
efficiency, enhanced reaction and sintering rate, cycle time and cost savings [6].
In the last four years, in this laboratory we have successfully sintered various
ceramics, composites, and even powdered metals to full density using microwave
processing [7,8]. Some highly transparent ceramic samples, such as alumina,
spinel, and aluminum nitride, have been successfully prepared by microwave
sintering process in our lab. Compared to the conventional sintering process, the
microwave sintering to highly transparent ceramic samples can be conducted at
lower sintering temperatures and much shorter sintering times [9].
EXPERIMENTAL
The ALON green samples in this work were prepared by mixing high purity
-Al2O3 powder (SM8, Baikalox, Baikoski International, NC, USA) and AlN
powder (Grade C, H.C. Starck, Laufenburg, Germany). The properties of the
starting powders are shown in Table 1. It was found that the addition of a small
amount of Y2O3 increased the densification and improved the transparency of the
sintered bodies during microwave sintering. Therefore the starting mixture
contained 67.5 mole percent of Al2O3, 33.5 mole percent of AlN, to which 0.5%
(by weight) Y2O3 in form of Y(NO3)3 6H2O was added. The powders with 3
wt.% of binder (Acryloid) were ball-milled in acetone for 24 hours. After drying,
the mixture was compacted uniaxially into pellets of diameter 12.7 mm and height
3 mm at a pressure of 30 MPa. Finally, the pellets were cold isostatically pressed
at 250 MPa for 5 minutes.
Table 1. The physical properties of the starting powders.
Manufacturer Purity Particle size Main phase
AlN powder H.C Starck (Grade C) >98% 2.41 m AlN
Al2O3 Powder Baikowski (SM8) 99.99% 0.15 m -Al2O3
Microwave sintering was carried out by using a 1.5 kW, 2.45 GHz single
mode microwave applicator in flowing pressure nitrogen at ambient pressure.
The heating rate was kept around 100 C/min by controlling the incident
microwave power. The phase composition of the samples was determined by x-
ray diffractometry (XRD). The densities of the sintered ALON samples were
measured by the Archimedes method. The optical microscope (Olympus) was
588 Ceramic Armor Materials by Design
used to study the microstructures, and the Varian spectrophotometer (CARY
2300) was used to measure the transmittance of the sintered samples.
RESULTS AND DISCUSSION
Figure 1 shows the XRD patterns of the starting mixture and microwave
processed samples under different synthesis conditions. The phase composition of
the starting material is pure -Al2O3 and AlN. The ALON phase appeared when
the sample was microwave heated at 1650ºC for only 10 minutes. The content of
ALON phase increased with the firing time at that temperature. A single phase
ALON was found after microwave firing at 1650ºC for 60 minutes.
Figure 1. The X-ray diffraction patterns of the starting mixture and
microwave synthesized ALON samples.
(a) Starting Material;
(b) Microwave synthesized at 1650°C for 10 minute;
(c) Microwave synthesized at 1650°C for 60 minute;
Figure 2 shows the densification behaviors of the ALON samples during
microwave sintering process. All microwave sintered samples exhibited only a
pure ALON phase composition which was confirmed by XRD. The theoretical
density (T.D.) of ALON is around 3.67 g/cm3. It was found that the samples
sintered at 1700 C for 1 hour with the density of 3.60 g/cm3 (~98.1% T.D.) were
still opaque or very slightly transparent. The samples sintered at 1750 C for 1
hour with the density of 3.67 g/cm3 (~100% T.D.) were quite translucent.
Ceramic Armor Materials by Design 589
However, by raising the sintering temperature to 1800 C and keeping the dwell
time unchanged, the grain size increased, and the transparency of the samples was
greatly improved. We tried to microwave sinter ALON at 1850 C, but the
sintering process was unstable because sometimes a discharge occurred which
resulted into partial melting of the samples. At the sintering temperature of
1800 C, the density of the samples increased with the increasing dwell time and
the transparency improved as well.
10 20 30 40 50 60
3.55
3.60
3.65
3.70
Sintering temperature = 1800o
C
Den
sity, g
/cm
3
Sintering time, min.1700 1750 1800
3.55
3.60
3.65
3.70
Sintering time = 60 min.
Density, g/c
m3
Sintering temperature,o
C
(a) (b)
Figure 2. Densification behavior of the ALON samples during microwave
sintering with (a) temperature, and (b) time.
The microstructural developments of the ALON samples during microwave
sintering are shown in Figure 3. As mentioned above, the samples sintered at
1750 C and 1800 C both for 1 hour showed the same density. But the grain size
of the 1800 C sintered sample was around 40-50 m, much higher than that of the
1750 C sintered samples (around 10-20 m), and also the grain boundaries
became narrower and cleaner, total grain boundary volume also reduced
dramatically. This obviously resulted in a transparency improvement. It was very
difficult to find pores in the sample sintered at 1700 C for 1 hour, that means the
sample had had a good densification, but the grain size was very small (less than
1-2 m), with very high grain boundary volume.
590 Ceramic Armor Materials by Design
(a) (b) (c)
Figure 3. Microstructures of microwave sintered ALON samples at (a)
1700 C, (b) 1750 C and (c) 1800 C for 1 hour.
Compared to single crystals, sintered polycrystalline ceramics (such as
ALON) have much more complicated microstructures including grains, grain
boundaries, second phases and pores. A light incident to a sintered body
experiences a diffuse reflection at the surface, and is subsequently absorbed and
scattered by the inhomogeneities inside the sintered body. To increase the
transmissivity of a sintered polycrystalline ceramic body, it is very important to
reduce porosity and the grain boundary phases since they strongly scatter light.
The ALON sample sintered at 1700 C for 1 hour had a high density up to 98%
T.D., but the grain structure was not developed well enough, and the grain
boundary volume was too large to cause considerable scattering of light. This
made the sample opaque.
0.5 1.0 1.5 2.0 2.5
0
10
20
30
40
50
60
70
80
90
100
Tra
nsm
itta
nce
, %
Wavelength, m
Figure 4. Transmittance of the ALON sample made by microwave sintering
at 1800 C for 1 hour (sample thickness = 0.6 mm).
Ceramic Armor Materials by Design 591
Figure 5. Appearance of the ALON sample made by microwave sintering at
1800 C for 1 hour.
Figure 4 shows the transmittance data of the microwave sintered ALON at
1800C for 1 hour. The total transmission of 60% was achieved for the polished
sample with a thickness of 0.6 mm. The sample shown in Figure 5 was optically
transparent. The results shown in this work have demonstrated that the
microwave sintering process can offer lower sintering temperature and much
shorter sintering times, in comparison with the conventional sintering process, to
make fully transparent ALON ceramics.
CONCLUSION
Fully transparent aluminum oxynitride (ALON) ceramic was successfully
prepared by pressureless microwave sintering processing. It was found that single
ALON phase formed at 1650 C in 60 minutes during microwave processing, and
the fully dense and highly transparent ALON samples were made at 1800 C with
residence time of 60 minutes.
ACKNOWLEDGMENTS
This work is partially funded by by DARPA /ONR under Grant No.
N00014-98-1-0752.
REFERENCES
1. T.M. Hartnett, S.D. Bernsein, E.A. Maguire, and R.W. Tustison, Optical
Properties of ALON (aluminum oxynitride), in Window and Dome
Technologies and Materials V, Proceedings of SPIE, edited by R.W.
Tustison, Vol. 3060 (1997)
2. N.D. Corbin, Aluminum Oxynitride Spinel: A Review, J. Euro. Ceram.
Soc., 5, 143-154 (1989)
3. R.L. Gentilman, E.A. Maguire, and L.E. Dolhert, Transparent Aluminum
Oxynitride and Method of Manufacture, US patent, 4720362 (1988)
592 Ceramic Armor Materials by Design
4. H.X. Willems, M.M.R.M. Hendrix, G. de With, and R. Metsalaar,
Production of Translucent –Aluminum Oxynitride, in Euro-Ceramics II,
edited by G.Ziegler and H. Hausner, Vol.3, 2443-2447 (1991)
5. J.W. McCauley, and N.D. Corbin, Phase Relations and Reaction Sintering
of Transparent Cubic Aluminum Oxynitride Spinel (AlON), J. Amer. Cer.
Soc. 62, 476-479 (1979)
6. W. H. Suttoon, in Microwave Processing of Materials III (R. L. Beatty,
W. H. Sutton, and M. F. Iskander, eds), Proceedings of the Materials
Research Society, Vol. 269, pp. 3-20 (1992)
7. R. Roy, D. Agrawal, J. Cheng, and M. Mathis, in Microwave: Theory
and Application in Materials Processing IV, Ceramic Trans., Vol. 80,
3-26 (1997)
8. R. Roy, D. Agrawal, and J. Cheng, Microwave Electromagnetic
Processing Invades New Materials, presented at the 2nd
World Congress
on Microwave & Radio Frequency Processing, Orlando, FL, USA, April
2-6, (2000)
9. J. Cheng, D. Agrawal, Y. Zhang, B. Drawl, and R. Roy, Fabrication of
Transparent Ceramics by Microwave Sintering, American Ceramic Society
Bulletin, Vol. 79, No. 9, 71-74, Sept. (2000)
Ceramic Armor Materials by Design 593
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AN INVESTIGATION OF THE TRANSMISSION PROPERTIES AND
BALLISTIC PERFORMANCE OF HOT PRESSED SPINEL
Mark C.L.Patterson
Technology Assessment & Transfer Inc.,
133, Defense Highway
Annapolis, MD 21401
Don W. Roy
21210 North 132 Drive,
Sun City West, AZ 85375
Gary Gilde
US Army Research Laboratory
AMSRL-WM-MC Building 4600
Aberdeen Proving Grounds
Aberdeen, MD 21005
ABSTRACT
The fabrication of transparent polycrystalline spinel (MgAl2O4) is being
pursued by Technology Assessment and Transfer Inc. (TA&T), with the goal of
becoming a commercial producer of transparent armor as well as optical windows
and domes. The process is based on hot pressing followed by hot isostatic
pressing to further improve the optical properties. This approach promises to
produce spinel at a cost significantly less than sapphire or ALON and at a scale up
to 22” diameter in the near term and possibly up to 36” in diameter. The larger
plates should be possible if the hot isostatic pressing step can be eliminated. This
paper discusses the effort underway to improve the optical properties of spinel
during hot pressing alone thereby establishing a low cost approach for transparent
armor.
The ballistic performance of spinel has been evaluated against ALON and
sapphire and the key properties of spinel are discussed with reference to its use in
infra-red windows and domes. High transmission in the mid infra-red is driving
a renewed interest in spinel for many optical systems. This paper provides an
overview of a joint effort between the Army and Technology Assessment and
Transfer Inc. to establish a capability for large spinel plate fabrication, and of
efforts to improve the optical transmission for multimode window and dome
applications.
Ceramic Armor Materials by Design 595
To the extent authorized under the laws of the United States of America, all copyright interests in this publication are the propertyof The American Ceramic Society. Any duplication, reproduction, or republication of this publication or any part thereof, withoutthe express written consent of The American Ceramic Society or fee paid to the Copyright Clearance Center, is prohibited.
INTRODUCTION
Magnesium aluminate spinel (MgAl2O4), a cubic oxide ceramic, has been
successfully sintered from selected reactive powders to transparency in the 0.3 m
to 5.5 m range. Transparency was first demonstrated in 1961 by the General
Electric Company and since that time there has been an intermittent effort to
develop optical quality spinel for a range of IR window, dome and armor
applications. A good summary of the early development efforts can be found in
the literature1 and in earlier work by the present authors2.
There are two main approaches to the fabrication of transparent spinel; the
first is by hot-pressing (HP) to transparency followed by hot isostatic pressing
(HIP) and the second is by pressureless sintering to produce an opaque product
which can be HIPed to transparency. Using HP/HIP processing, excellent optical
performance was achieved previously with spinel in small sizes and thin wall
thickness by Coors Ceramics and Alpha Optical Systems. Spinel domes were
qualified for at least two IR guided missiles and for the stinger missile launch tube
window, prior to the shutdown of production when military budgets were reduced
following the Gulf War.
An effort to fabricate transparent spinel by RCS Technologies Inc., using “rate
controlled pressureless sintering” followed by HIP processing in the early 90’s
showed considerable promise, but could not be sustained because of the lack of
financial support3. Consequently there has been no commercial spinel production
since 1993.
There is a need within the military to reduce the weight and increase the size
capability of transparent armor systems while simultaneously increasing ballistic
protection capabilities. Additionally, there is growing need for window and
dome materials which extend further into the IR and can be used for multimode
weapons systems that are exposed to very demanding environments.
Polycrystalline MgAl2O4 spinel has been recognized for many years as a material
with great potential for transparent armor and for UV, visible and mid IR optical
component applications. Based on these needs, TA&T is scaling up to produce
spinel commercially for both optical and armor applications
This work is being driven not only by the need for spinel based on its unique
properties but based on current as well as prior manufacturing information. It is
expected that the cost of large spinel plates will be significantly less than the
competitive ALON and sapphire materials.
Properties of Spinel
Spinel crystal structure is cubic and optically isotropic; thus polycystalline
shapes may be fabricated without severe scattering problems inherent in
596 Ceramic Armor Materials by Design
polycrystalline non-cubic materials. In the microwave region the isotropy of
spinel prevents localized absorption and heating that occurs in non-cubic
materials because of differing grain boundary orientation and anisotropic
dielectric loss index. Spinel undergoes no polymorphic transformations, so it is
free of problems due to thermally induced phase changes. Extensive programs
were carried on in the 1980’s to measure the properties of spinel as well as other
candidate window materials, including sapphire, ALON and yttria at Johns
Hopkins University Applied Physics Laboratory4 and Honeywell Systems
Research Center5. The typical physical properties for polycrystalline spinel are
listed in Table I.
Table I Typical Physical Properties of Polycrystalline MgAl2O4 Spinel
Density 3.58 gm/cc
Hardness, Knoop [100gm] 1398kg/mm2
Minimum Strength @ 25oC
4-pt bending 15x 103psi [103x10
6Pa]
Biaxial 25x 103psi [172x10
6 Pa]
Tension 16x103psi [110x10
6 Pa]
Compression
390x103psi[2.69x10
9Pa]
Elastic Mod. 39x106psi 273x10
9Pa]
Bulk Mod. 27.9x106psi 192x10
9Pa]
Shear Modulus 15.9x106psi[110x10
9Pa]
Thermal Coefficient of Expansion
25 - 200oC 5.6x10
-6/oC
25 - 5000C 7.3x10
-6/C
25 - 1000oC 7.9x10
-6/C
Specific Heat, cal/gm/oC
20oC 0.21
Poisson’s Ratio 0 .26
Dielectric Strength, kV/mm
.05’’[1.27mm] thick 490
.01’’[.25mm] thick 580
Melting Point 21350C [ 3875
oF]
Volume Resistivity, ohm-cm
25oC >10
14
300oC 5x10
14
500oC 2x10
11
700oC 4x10
8
Thermal Conductivity,
gm-cal/cm2/sec/
oC [W/m-
oK]
25oC 0.060 [24.7]
100oC 0.0357 [14.8]
1200oC 0.0130 [5.4]
Dielectric Constant & Loss index
1KHz 8.2 0.00025
1MHz 8.2 0.0002
9.3GHz 8.3 0.0001
The refractive index of spinel has been measured to vary between approximately
1.74 and 1.66 over the range of its transparency as shown in Table II.
Table II. Reflective index of spinel at different wavelengths.
( m) 0.49 0.59 0.66 1.0 2.0 3.0 4.0 5.0
Ref.Ind 1.736 1.727 1.724 1.704 1.702 1.698 1.685 1.659
Spinel has distinct optical property advantages over both sapphire and ALON.
In contrast to cubic spinel, single crystal Al2O3 [sapphire] is anisotropic and
birefringent, causing optical design problems11
. ALON has a shorter transmission
cut-off in the 4.5 to 5.5 micron spectral region, resulting in a significantly higher
Ceramic Armor Materials by Design 597
coefficient of absorption in that critical band. This is shown dramatically in Table
III where the relative transmission properties at RT, 250 C and 500 C are shown
for spinel, sapphire and ALON. At 4.8 microns and 250 C spinel offers a 4% and
14% improvement in transmission over sapphire and ALON respectively.
Table III. Transmission properties for spinel, sapphire and ALON vs. wavelength
at temperatures up to 500 C6.
Percent Transmission at Wavelength in ( m)
Sample & Temp.( C) 3.0 4.0 4.5 5.0 5.5 6.0
25 88 87 77 59 11 -
ALON 250 87 84 71 46 7 -
500 87 81 62 33 2 -
25 87 87 84 71 49 -
Sapphire 250 86 86 79 61 32 -
500 84 82 72 50 20 -
25 84 87 84 76 54 22
Spinel 250 84 87 82 67 39 11
500 82 83 76 55 23 4
Spinel Applications
As an optical material spinel is similar in nature to both ALON and sapphire
in that it has a high hardness, erosion resistance, and transmits from
approximately 0.25 m to 6.0 m. It is isotropic and does not therefore exhibit the
birefringence seen in sapphire and as discussed above exhibits a lower absorption
coefficient than either sapphire or ALON in the mid infrared, particularly at
elevated temperatures.
Based on these properties and the hope that spinel can be fabricated at a
considerably reduced cost over either ALON or sapphire, spinel is being
developed for use as erosion resistant multimode windows and domes for a wide
range of defense applications. It is also being investigated for optical lenses as
well as armor against hard, armored piercing projectiles. For armor applications,
the possibility of producing large panels, possibly up to 36” in diameter through a
low cost hot-pressing process is driving the continued interest at present.
Ballistic Evaluation of Spinel
The Army has been interested in spinel for transparent armor since the late
sixties7. When all factors including transparency, hardness, impact resistance,
strength, modulus, ease of fabrication, and crystal size capability are taken into
consideration, spinel appears particularly well suited for armor applications.
Because of the U.S. Army's continued interest in spinel for transparent armor and
598 Ceramic Armor Materials by Design
electromagnetic windows, ARL and TA&T signed a cooperative Research and
Development Agreement (CRADA) in 1998 for the development and dual use
assessment of transparent spinel, using hot-press/HIP processing. Good
transparency has been achieved in flat plates up to 5 inches square and 0.5 inches
thick. While the main focus of recent work has been for the fabrication of thick
section (>0.4 inches) spinel plate, efforts have also included the evaluation of
multiple pieces of thinner sectioned plates.
A recent evaluation of transparent spinel and ALON carried out by the ARL8
demonstrated that both materials dramatically improve the performance of
transparent armor systems over the traditional glass/plastic systems currently in
use, based on areal density and velocity requirements9. The results of this and
other ceramic based systems compiled between 1969 and 1996 are shown in
Figure 1. For this particular threat, the data shows that spinel backed with
polycarbonate performs approximately 4% better than ALON and 10% better than
sapphire, both backed with polycarbonate. The spinel/polycarbonate was
approximately 1.5 lbs/ft2 heavier than the spinel/polycarbonate system. With an
areal density of 12 lbs/ft2 it exhibited a V50 between 2900 ft/sec and 2950 ft/sec.
ALON backed with both glass and polycarbonate exhibited a V50 of 3000 ft/sec.
The weight of the armor system, however, was 33% higher at 16 lbs/ft2.
- Spinel
- ALON
- Sapphire
- Ball. glass
V50
Inverse areal density
Desired ballistic
performance
Figure 1. Relative ballistic data for spinel, sapphire and AlON against an
unspecified threat.
There is presently a growing need for lightweight transparent armor concepts
against armored piercing (AP) 12.7mm projectiles, which can fulfill the mass
and/or thickness (scale) requirements for air and light armored vehicle programs.
Spinel in thin layers has in the past been laminated with glass and polycarbonate
backing to defeat 7.62mm AP projectiles. Recently, thicker sections of spinel (up
Ceramic Armor Materials by Design 599
to 20mm) have been evaluated for this application10 and its performance
compared with Al2O3 and SiC against 12.7mm projectiles at two different
velocities. The spinel tested for this application was fabricated in tiles up to 20cm
square by the French company, Ceramiques & Composites. The ceramic front
face was laterally confined and bonded to an aluminum honeycomb back surface
(no ballistic influence). Ballistic evaluation was carried out at 2880 ft/sec and
1800 ft/sec. In their study they determined that spinel outperformed alumina at
both projectile velocities but was inferior to SiC. (The alumina was not sapphire
but was a 92% Al2O3 ballistic grade). They also determined that ballistic
protection against AP 12.7mm at 2880 ft/sec could be obtained from a sapphire
front surface backed with polycarbonate at an areal density of approximately 21.5
PSF or by a glass/polycarbonate laminate with an areal density of approximately
41.0 PSF. They estimated that the same protection with spinel could be achieved
with an areal density of 20.5 PSF.
Recent evaluation of spinel tiles tested at the Army Research Laboratory in
Aberdeen showed comparable performance for both spinel and ALON, and a
significant weight reduction over a glass system. The tiles were 4” square by
0.375” thick and shot with a steel core, small caliber projectile. The actual data is
not available but is shown normalized in Table IV together with a baseline
glass/plastic system.
Table IV. Normalized ballistic performance (V50 and areal density) for spinel
ALON and a glass/plastic baseline against a small caliber, steel cored projectile.
Glass/plastic Spinel ALON
Areal Density 1.0 0.43 0.44
V50 1.0 0.88 0.89
Although ALON and sapphire are very promising transparent armor materials,
spinel may be able to offer the best balance of both performance and affordability.
PROCESS DEVELOPMENT
Technology Assessment and Transfer Inc., is pursuing a HP process using LiF
as a sintering aid. At present the process requires subsequent HIP to produce
good optical transmission with low haze, but the goal of future development will
be to investigate if the final HIP step can be eliminated, thereby significantly
reducing the cost and allowing the fabrication of parts up to 36” in diameter.
The theory behind the present pressing approach is to ensure that volatile
contaminants such as LiF are allowed to escape from the sintering body before
closed porosity is attained. If the grains are allowed to sinter at too high a
temperature bridges are formed, which upon application of further load cannot be
600 Ceramic Armor Materials by Design
broken, resulting in opacity. It is important therefore to balance application of the
load with out gassing of the sintering body and temperature to ensure continuous,
yet gradual microstructural development. To date, satisfactory pressing
procedures have been established to produce good transmission (>80%) in 0.40”
thick sections of spinel up to approximately 5” in diameter. The present focus of
work is threefold:
understanding the effect of processing environment on microstructural
development and properties
optimize pressing procedures for shaped configurations including domes
increase the size capability to 22” in diameter by year end 2002
Processing, Microstructural Development and Properties
It has been shown previously that the transmission properties of spinel can be
improved if the temperature is increased, annealed11, or if it is HIPed12. Figure 2
shows the increase in transmission that has been observed when samples are
annealed for longer periods of time at a temperature below the sintering
temperature.
Figure 2. Percent transmission up to 3 microns following HP at 1650 C for 3 hours
(right) and following an additional anneal at 1550 C for 12 hours (left).
Initially it was proposed that the increase in transmission following HIPing
was attributed to a growth in the grain size and a reduction in the number of grain
boundaries. A complimentary increase in the strength following HIPing, which
was also observed previously by Don Roy, was thought to be due to grain
boundary development (a reduction in the impurity levels or realignment of
adjacent grains). A brief study was undertaken to investigate the microstructural
development that takes place during HP, annealing and HIP. Some of theses
results are described below.
Ceramic Armor Materials by Design 601
Orientation Imaging Microscopy: Microstructural analysis was performed on
three samples following HP at 1650 C, HP at 1650 C followed by annealing at
1550 C for 12 hours, and following HP at 1650 C and HIP at 1700 C. Once
polished, each of these samples showed transmission values summarized in Table
V showing the significant increase in transmission that can be achieved in the
visible and near UV regions. In the near IR regions the increase in transmission is
less pronounced. It was hoped that the cause for this increase in the optical
transmission could be seen in the microstructure and so orientation imaging
microscopy was performed on these samples using a FEI XL-30 FEG SEM with a
beam current of 1.5nA. Data was collected from an area 1,200 m by 2,400 m in
size using a step size of 5 m.
Table V. Transmission and microstructural data for spinel following HP -
1650 C, HP & annealing at 1550 C, and following HP & HIP - 1700 C.
Attribute HP only HP/anneal HP/HIP
Wavelength 0.5 m 2.0 m 0.5 m 2.0 m 0.5 m 2.0 m
Transmission 45% 70% 65% 78% 78% 79%
Av grain size 38.3 m 39.5 m 45.1 m
Aspect ratioa
0.55 0.58 0.65
Av.misorientationb
~ 45 ~ 45 ~ 45
Textc. all grains 2.678 2.927 3.329
Textc. small 2.061 2.111 2.365
Textc. medium 2.556 2.328 2.836
Textc. large 6.864 6.000 7.405
In an effort to establish whether or not grain texture was contributing to the
overall transmission properties of the spinel, density pole figure plots were
generated for each of the samples for the 100, 101 and 111 axis. These were
generated for all the grains as well as for individual groups of grains as described
in Table V above. The maximum values measured for any of the reference
orientations are shown plotted as a times random (1.0) value in Table V, showing
that there is slightly more texture seen in the HIP spinel and that texture is driven
primarily by the large grains. The average grain size is similar for the HP and HP
annealed sample but is larger in the HP/HIP sample, as shown in Table V. The
a grain aspect ratio compared with 1.0 being equiaxed.
b av. misorientation angle distribution for all grains is estimated as being the same
c maximum times random values (1.0) from summary of pole figure plots. Small refers to grain
less than 50 m, medium 50-120 m and large above 120 m in diameter.
602 Ceramic Armor Materials by Design
grain aspect ratio increased for the annealed and the HIPed samples. However, the
misorientation angle distribution was centered around 45 for all the samples.
Figure 3. Image quality maps showing the grain size distributions for the samples
described in Table V. The shaded regions in the distributions refer to bands of
50 m to 100 m and 300 m to 500 m respectively.
Figure 3 shows the grain boundary structure for each of the three samples and
highlights very low angle grain boundaries (1 -5 ), low-angle grain boundaries
(5 -15 ), and high angle grain boundaries (>15 ).
Ceramic Armor Materials by Design 603
Although it may not be clear in these micrographs, it was interesting to note
that the very low angle grain boundaries (as indicated by the arrows) seemed
concentrated within certain grains and were not evenly distributed throughout the
microstructure in any of the samples. No orientation correlation was apparent
with any of the selected grains.
Results of the Orientation Imaging Microscopy: This initial analysis of the spinel
microstructures following different treatments revealed little insight as to the
significantly higher transmission values which were observed following annealing
of the samples. Following HP/HIP it can be seen that there is a larger average
grain size, a shift in the grain size distribution towards a higher fraction of large
grains, a higher fraction of near equiaxed grains, and a higher fraction of
directional alignment. These differences are not observed between the HP and
HP/annealed spinel samples. It is interesting to note that no correlation could be
seen in the distribution of very low angle grain boundaries, which appeared
localized within specific grains and not evenly distributed throughout the
microstructure of all three spinel types. Additionally, no correlation could be seen
between these grains and their orientation.
The increased transmission values were most pronounced in the near UV as
shown in Figure 2, and it is expected that the influences may occur at a smaller
scale than were observed and evaluated in this study.
PRODUCTION APPROACH
A facility is presently being made ready through the installation of processing
and quality assessment equipment. The first of these to be installed into the new
facility is a 600 ton Birdsboro press with 72” of daylight. A heating package and
vacuum enclosure will be installed separately, thereby allowing the fabrication of
plates up to 22” in diameter (and possibly up to 36” in diameter in the future).
The 600 ton Birdsboro press is shown partially constructed in Figure 4. Each
of the four posts (not shown in this picture) are 12” in diameter and the main ram
is 16” in diameter. It is expected that through further changes to the hydraulics
system it will be possible to increase the load capability to 1000 tons, thereby
allowing even larger spinel plates to be fabricated. In addition to fabrication of
single components with large diameters, this press will also be used to process
multiple parts in the form of domes, lenses or smaller windows.
604 Ceramic Armor Materials by Design
Figure 4. 600 ton press during construction following alterations
The spinel fabrication process that has been selected by TA&T is hot-
pressing, followed by hot isostatic pressing. The initial hot-pressing process uses
LiF as a sintering aid and results in a transparent product which can readily be
inspected for internal flaws, inclusions and discoloration. The second hot
isostatic pressing step further improves the optical properties of the spinel and
reduces variation in other mechanical properties such as strength. A process flow
chart for the individual operations is shown in Figure 5.
The hot-pressing process is a forgiving one in which good optical properties
can be obtained within a single processing step. At this time the optical properties
required for window and dome applications can only be achieved through both
hot-pressing and hot isostatic pressing. It is expected that with future
development it maybe possible to achieve the required optical properties from a
single hot-pressing step, thereby leading to a significant reduction in the
processing costs. This approach will be investigated over the next 2 years. As yet,
a two step process is still required. To date the hot-pressing process has been
Ceramic Armor Materials by Design 605
capable of producing flat plates of spinel up to a thickness of approximately 0.5”
and has also shown the ability to fabricate near hemispherical domes
approximately 6” in diameter. Additionally, the hot-pressing approach has
historically provided a high yield in excess of 60% as compared with alternative
approaches such as pressureless sintering13.
RejectReject
Final finishing
processes
Hot, isostatic
pressing of partQuality
controlQuality
control
Hot-press
spinel part
Binder
burnout
Formation of
“green” shape
Powder blending
and mixing
Figure 5. Process flow chart showing the individual operations required for
optical spinel fabrication
DISCUSSION
Technology Assessment and Transfer Inc. is establishing a facility to become
a commercial supplier of transparent spinel parts for armor, IR window and dome
applications. The focus of the present work is to understand the effect of
processing environment on microstructural development and properties of the
spinel. The pressing procedures are being optimized for flat plates and shaped
configurations including domes, and the size capability is being increased up to
22” in diameter by year end 2002.
An improvement in the optical transmission is seen following HP if the spinel
is annealed or HIPed. Following HP/HIP it can be seen that there is a larger
average grain size, a shift in the grain size distribution towards a higher fraction
of large grains, a higher fraction of near equiaxed grains, and a higher fraction of
directional alignment. However, following annealing, these differences were not
observed and the cause of the improved optical performance over the HP spinel is
as yet unknown.
The HP processing of spinel promises to be a low cost approach to producing
transparent armor with similar performance to both ALON and sapphire but with
the capability of producing larger panels than are possible with either ALON or
sapphire. In the near term 22” diameter plates will be made and this will possibly
be increased to 36” in the future. Microstructural development of spinel is
presently being investigated with the purpose of improving the optical properties
606 Ceramic Armor Materials by Design
of spinel without the HIP requirement, thereby reducing significantly the overall
fabrication costs. The HP process produces high yields and is capable of
fabricating near net shape parts, such as full hemispherical domes. Greater than
80% transmission has presently been achieved in 0.4” thick parts 5” in diameter
and a number of parts formed from near net-shape HP/HIP are presently being
evaluated for optical application.
ACKNOWLEDGEMENTS
The authors would like to thank Matt Nowell at TexSEM Laboratories for
performing orientation imaging analysis on the spinel samples. This work was
funded in part by the Army Research Laboratory Aberdeen under an SBIR
contract # DAAD17-00-C-0080.
REFERENCES
1 W.H.Rhodes, “Phase Chemistry in the Development of Transparent
Polycrystalline Oxides”, GTE Laboratories, TR-0209-07-92-082, 1992. 2 D.W.Roy, M.C.L.Patterson, J.E.Caiazza and G.Gilde, “Progress in the
Development of Large Transparent Spinel Plates”, 8th
DoD Electromagnetic
Windows Symposium Proceedings, ASAFA Colorado Springs, CO 24th
-27th
April
2000.3 M.L.Huckabee, “ Near net shape spinel optics for broadband windows,
lenses and domes” Final report contract DAAH04-95-C-0010, RCS Technologies
Inc. 1995 4 M.E.Thomas, R.L.Joseph and W.J.Tropf, “Infrared Properties of Sapphire,
Spinel and Yttria as a Function of Temperature”, SPIE vol. 683, 1986. 5 J.A.Cox, D.Greenlaw, G.Terry, K.McHenry and L.Fielder., “Infrared and
Optical Transmitting Materials”, SPIE vol. 683, 1986. 6 SWIR/LWIR “Optical Sensor Window Development Program”. Final
Report DASG60-85-C-0018. 7 A. Gatti and J. Noone, Feasibility Study for Producing Transparent Spinel,
General Electric Company, Space Sciences Laboratory , Space Division, King of
Prussia, PA Final Report for Contract DAAG46-69-C-0096 8 M.C.L.Patterson, J.E.Caiazza, and D.W.Roy, “Transparent Spinel
Development”, Inorganic Optical Materials II, Alexander J.Marker III, Eugene G.
Arthurs Editors, Proc, of SPIE Vol. 4102 pp.59-68. 2000. 9 J.Conners., “Magnesium Aluminate Spinel, Material and Prototype
Development,” ARL Internal Communication, 17th June 1997. (This article was
not seen by the author but has been quoted directly from a paper by J.J.Swab,
J.C.LaSalvia, G.A.Gilde, P.J.Patel and M.J.Motyka, “Transparent Armor
Ceramics:ALON and Spinel”.
Ceramic Armor Materials by Design 607
10 C.E. Cottenot, “Transparent Ceramic for Lightweight Armors,” Lightweight
Armor Systems Symposium ‘95 Cranfield, England, 28-30th June, 1995. 11
M.C.L.Patterson, G.Gilde and D.W.Roy, “Fabrication of Thick Panels of
Transparent Spinel” Inter. Symp. Proc. Optical Science & Technology. SPIE 46th
Annual Meeting San Diego, CA. 29th
July – 3rd
August 2001. 12
G.Gilde, P.Patel and M.C.L.Patterson, “A comparison of hot-pressing, rate
controlled sintering and microwave sintering of magnesium aluminate spinel for
optical applications”, SPIE Conf on Window and Dome Technologies and
Materials VI, Orlando FL. April 1999. Vol. 3705. pp. 94-104. 13
Private communication, Don Roy August 2001.
608 Ceramic Armor Materials by Design
Microstructure and Macrostructure Effects
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THE EFFECT OF MICROSTRUCTURE ON THE DYNAMIC BEHAVIOR OF
COMPOSITE ALUMINA/TITANIUM DIBORIDE
Kathryn V. Logan, Ph.D., P.E.
School of Materials Science and Engineering
Georgia Institute of Technology
Atlanta, Georgia 30332-0245
ABSTRACT
Past work has shown that the dynamic behavior of a dense, hot pressed
ceramic-ceramic composite that is composed of a nominal 70wt% alumina/
30wt% titanium diboride formed from powders that were either manually mixed,
or synthesized using self-propagating high temperature synthesis (SHS) is
significantly affected by the microstructural bias, including phase distribution and
grain morphology, formed during synthesis and processing. A method has been
developed to bias the composite microstructure such that the titanium diboride
grains are caused either to be dispersed amongst, or to surround the alumina
grains. A review of past work on the significance of processing/forming
consistency, as well as results to date on efforts towards quantitative
characterization of the microstructure are presented.
INTRODUCTION AND REVIEW OF PAST WORK
A number of ceramic materials having potential application as high strain rate
armor materials have exhibited superior ballistic mass efficiencies comparable to
steel, but the results have not been consistent and the material properties that induce
a resistance to high strain rate penetration have not been completely determined,
especially for the effects of various microstructures.
In order to elucidate some of the mechanisms in ceramics that optimize armor
performance, a systematic study on controlled microstructure alumina/titanium
diboride ceramic composite materials has been carried out over the last several
years. Long rod penetrator (LRP), depth of penetration (DOP) tests conducted at
Aberdeen Proving Ground (APG) on these materials exhibited superior performance
with ballistic mass efficiencies up to four.1 The multi-phase material was a dense
ceramic particulate composite with a preferentially biased microstructure of hard
TiB2 grains congregating around grains of less hard Al2O3.
Ceramic Armor Materials by Design 611
To the extent authorized under the laws of the United States of America, all copyright interests in this publication are the propertyof The American Ceramic Society. Any duplication, reproduction, or republication of this publication or any part thereof, withoutthe express written consent of The American Ceramic Society or fee paid to the Copyright Clearance Center, is prohibited.
Previous work by Logan2 demonstrated the ability to influence the
microstructural bias of a hot pressed alumina/titanium diboride composite, both in
starting powders that were either produced using Self-Propagating High
Temperature Synthesis (SHS) or commercially available powders that were
manually mixed (MM). The microstructure of the composites can be preferentially
biased such that the morphology and distribution of the component phase grains can
be partially controlled.3,4
The microstructure designated as “A” (figure 1)
represents SHS powders hot pressed to >98% theoretical density ( t) forming a
microstructure that is biased towards 1-10 m titanium diboride grains (white
areas) comprising an average 7.4 m phase size surrounding 20-40 m alumina
(gray areas) grains. The microstructure designated as “B” (figure 2) represents
SHS powders hot pressed to >95% t forming a microstructure that is biased
towards 1-5 m titanium diboride grains (white areas) comprising an average 6.2
m titanium diboride phase size that is uniformly distributed amongst 10-20 m
alumina (gray areas) grains. The microstructure designated as “C” (figure 3)
represents MM powders hot pressed to >98% t forming a microstructure that is
biased towards 1-10 m titanium diboride grains (white areas) comprising an
average 8.7 m titanium diboride phase surrounding an alumina phase (gray
areas) comprised of grains up to 100 m in diameter. The microstructure
designated as “D” (figure 4) represents MM powders hot pressed to >98% t
forming a microstructure that is biased towards 1-10 m titanium diboride grains
(white areas) uniformly distributed in an alumina phase (gray areas) comprised of
grains averaging 12.3 m.
Composites having the biased microstructures have exhibited quasi-static and
dynamic behaviors that indicate a tendency to vary because of the microstructural
bias. Keller and Zhou5 have found dynamic compressive strengths of the four
biased microstructures described above to range from 4.4 to 5.3 GPa; values
which are 27% higher than the quasi-static values. Also, the measured
compressive strength directly correlates with the fraction of titanium diboride rich
areas on the fracture surfaces. The failure associated with the alumina phase is
transgranular; while the failure associated with the titanium diboride phase is both
transgranular and intergranular. Table I is a summary of dynamic properties of
four representative biased microstructures A, B, C and D. Figures 5-10 are
fracture surfaces of SHS and MM composites after MOR bar breaks. Figure 5
(boxed area) shows cleavage in the relatively large alumina phases (gray areas)
with the titanium diboride (white areas) tending to be localized around large
(gray) areas of alumina (figures 5 and 7). Figures 6 and 8 show evidence of
titanium diboride grains being more homogeneously distributed in the alumina,
crack pinning by titanium diboride, and a high concentration of microcracks.
Figure 9 shows relatively large TiB2 grains (light gray) localized at the alumina
612 Ceramic Armor Materials by Design
grain boundaries (darker gray). Figure 10 shows the relatively large TiB2 grains
(white areas) homogenously distributed amongst the alumina grains (gray areas).
Figure 1. Sample A, (SHS T@A) Figure 2. Sample B, (SHS TinA)
[_____] = 100 microns
Figure 3. Sample C, (MM T@A) Figure 4. Sample D, (MM TinA)
Ceramic Armor Materials by Design 613
Table I. Summary of dynamic properties
Compressive
Strength
Spall
Strength Wave Speed
Hugoniot
Elastic Limit
Tensile
Yield
Strength
Sample (GPa @103 s-1) (GPa) (KM/S (GPa) (GPa)
A 5.2 0.32 8.24 +/- .83 6.2 +/- 3.1 4.2 +/- 2.1
B 4.6 N/A 9.67 +/- 1.0 4.4 +/- 1.2 3.11 +/- 0.84
C 4.4 0.311 9.08 +/- .74 5.5 +/- 2.3 4.02 +/- 1.7
D 5.3 0.222 8.31 +/- .78 8.5 +/- 4.5 6.23 +/- 3.3
TiB2 4.2 0.33 9-18
Al2O3 4.0 0.45 ~6.7
Figure 5. MOR bar fracture surface Figure 6. MOR bar fracture surface
Sample A, SHS (T@A) Sample B, SHS (TinA)
Kennedy, et. al.6 have determined that while the Hugoniot Elastic Limit and the
compressive strengths of the four biased microstructures are dependent on the
average polycrystalline grain (phase) size, the tensile (spall) strength scales with
the titanium diboride-phase connectivity. This result implies that the
interconnected microstructure provides a higher resistance to failure in tension
compared with a microstructure having homogeneously dispersed particles. Table
I is a summary of the observed dynamic properties. Table II is a summary of the
phase sizes.
614 Ceramic Armor Materials by Design
Figure 7. MOR bar fracture surface Figure 8. MOR bar fracture surface
m m
Sample A, SHS (T@A) Sample B, SHS (TinA)
Figure 9. MOR bar fracture surface Figure 10. MOR bar fracture surface
Sample C, MM (T@A) Sample D, MM (TinA)
Ferranti7 has found that processing parameters influence the development of
microstructural bias and composite properties in that the interconnected TiB2
polycrystalline phase forms directly during the SHS reaction. Mixing the B2O3
Ceramic Armor Materials by Design 615
with Al prior to combining with TiO2 promotes TiB2 phase connection following
the SHS reaction. Ball milling of the resultant SHS product reduces the inherent
phase connectivity with the phase size decreasing as milling periods are longer. A
bimodal particle size distribution of synthesized powders produced high-density
parts; hot-press parameters did not appear to affect TiB2 phase connectivity.
Table II. Summary of phase sizes.
Sample Average Integral TiB2 Phase Al2O3 Phase
Curvature ( m-1) Size ( m) Size ( m)
A -0.316 +/- 0.022 7.0 10.4
B -0.476 +/- 0.046 6.2 9.1
C -0.074 +/- 0.028 8.7 25.1
D -0.375 +/- 0.031 7.9 12.3
The properties of composite ceramics usually follow the Rule of Mixing and
are influenced by the properties of the continuous phase.8 However, the
processing mechanisms that allow control and reproducibility of a specific
microstructure in multi-phase ceramic materials is not completely understood; and
accordingly, how the processing factors would directly affect and optimize
properties in high performance structural applications. The potential high
performance properties of a material are defined by intrinsic properties such as
crystal structure, bond strength and composition. However, the actual material
performance is significantly influenced by extrinsic properties such as grain size,
porosity and phase distribution. Therefore, the ability to form materials having
superior performance properties requires a detailed knowledge and control of the
processing routes that influence microstructural development.
A number of processing parameters have been shown to influence the
resulting microstructure of the hot pressed SHS and MM alumina/titanium
diboride composites: the initial state of the hot pressed powders (stoichiometry,
particle morphology and proximity), and densification variables (pressure,
temperature, and time at temperature).
The compositions to date have been based on a (nominal) stoichiometry of
30wt% TiB2 and 70wt% Al2O3 according to equation (1)
3 TiO2 + 3 B2O3 + 10Al => 3TiB2 + 5Al2O3 (1)
(Note: The SHS reaction produced a product that was (adiabatically) 29% TiB2
and 71% Al2O3; while the MM powders were mixed to be 30% TiB2 and 70%
Al2O3.) It is probable that the compositions used were close to a eutectic or
616 Ceramic Armor Materials by Design
eutectoid composition since characteristic eutectic-like microstructures have been
observed in the SHS composites (figure 11).
Figure 11. Eutectic-like microstructure.
TiB2 (white areas) and Al2O3 (gray areas)
The crystal structure of titanium diboride is simple hexagonal close packed
(HCP) with a c/a lattice ratio of 1.07. The unit cell lattice parameters have been
reported as a = 3.03A, and c = 3.23A with the c/a lattice ratio remaining 1.07 from
25 C to 1200 C.9. The c/a lattice ratio of an ideal HCP unit cell is 1.63. Alpha-
alumina also has a hexagonal close packed crystal structure. The lattice parameters
are a = 4.76A, and c = 12.99A with a c/a ratio of 2.73. Grain morphology of a
polycrystalline, HCP crystalline ceramic material can vary in shape from equiaxed,
to hexagonal, to high aspect ratio lamellae. The number of faces and edges on a
grain, as well as the bond strength between the grains, will influence the ultimate
mechanical strength and behavior.
Prior observations have also shown that the microstructure of the SHS
composite titanium diboride is influenced by the rate of application of pressure to
the red hot, plastic product immediately after the reaction has occurred. A
comparison was made between the microstructure that was formed in-situ just
after the SHS reaction had occurred and the microstructures formed as pressure
was applied immediately after the reaction occurred when the product was still
red-hot and plastic. The applied pressures were 12.42MPa (1800psi), 17.25MPa
(2500psi), 20.70MPa (3000psi) and an explosively applied pressure. Figure 12 is
a micrograph of the microstructure formed in-situ after the SHS reaction has
occurred showing a heterogeneous localization of TiB2 (white area) in Al2O3
(gray area). The effect of application of pressure is shown in Figures 13-15: as
the rate of applied pressure is increased, the aspect ratio of titanium diboride
decreases.10
Ceramic Armor Materials by Design 617
Ceramic materials generally show some increase in compressive strength with
an increase in strain rate loading. The compressive strength of TiB2 increases with
application of pressure.11
and increases significantly with strain rate loading by
exhibiting a reported HEL of 160 kbar at a shock stress of 240 kbar.12
A marked
strengthening of alumina is shown with increasing strain rate.13
The strain rate
during chip formation in metal cutting is determined to be 104 s
-1 so the process
would be expected to approach adiabatic conditions.14
m
m
Figure 12. In-situ foam Figure 13. 12.42 MPa applied pressure
ELECTRICAL RESISTIVITY CHARACTERIZATION OF
MICROSTRUCTURES
Since the properties of alumina/titanium diboride materials are microstructure
dependent, it would therefore be advantageous if a simple non-destructive
measurement could be used to determine the degree of microstructural bias.
Titanium diboride (TiB2) is an intermetallic compound that has very impressive
performance characteristics:15,16,17
it acts very much like a metal with an electrical
resistance comparable to that of copper at 1000oC. The conductivity of TiB2 is
approximately 10-55 micro-ohm-cm at 300-1200K. As in a metal, the resistivity
of TiB2 decreases with increasing temperature.18
Therefore, the connectivity and
percolation path of the titanium diboride will govern the electrical properties. 19
Test bars were cut to dimensions specified in MIL-SPEC 1942B from each
hot pressed, three-inch OD disk of the four biased microstructures A, B, C and D.
The bars sampled from the discs were used to obtain mechanical and electrical
property data. The quasi-static and dynamic mechanical property results have
been previously reported,20
so only electrical property results21
will be reported
here. Because the surface conductivity of the composite was poor, silver paint
was used to create a good connection to the Fluke multi-meter leads. The
618 Ceramic Armor Materials by Design
m m
Figure 14. 20.70 MPa applied pressure Figure 15. Explosive pressure
resistance for each bar was recorded at two different times. Two readings were
taken to check the stability of the measurement and the consistency of the
measurement techniques.
Figure 16 is a summary of the electrical resistivity measurements. In general,
the average resistivity of the MOR bars that were cut from the hot pressed
manually mixed powders (0.213-4.137 ohm-cm) was lower than the average
resistivity of the MOR bars that were cut from the hot pressed SHS powders
(0.234-53.866 ohm-cm). The lower resistivity in the MM samples was probably
due to the relatively large titanium diboride grains as compared with the smaller
grains in the SHS sample. No significant affect on the resistivity could be
discerned between the composites with TiB2 segregated at the alumina grain
boundaries (0.213-0.336 ohm-cm) and the composites with TiB2 uniformly
distributed in alumina (0.234-0.266 ohm-cm). It was also found that a four hour
hot pressing hold time at temperature reduced the resistivity of both powder types:
MM (0.21-0.27 ohm-cm), SHS (0.23-0.40 ohm-cm). After the four-hour hold at
temperature, the SHS composite resistivity was comparable to that of the MM
composites. It was apparent that the TiB2 provided a path allowing reduction of
the overall resistivity.
SUMMARY
A. Processing variables significantly affect resulting microstructure and thus
performance of hot pressed SHS and MM powders.
B. Although Samples A, B, C and D showed tendencies towards a preferentially
biased microstructure allowing one to discern trends in microstructural effects on
Ceramic Armor Materials by Design 619
performance, further research is necessary to determine the specific processing
parameters to produce a totally biased microstructure.
C. The dynamic behavior of composite alumina/titanium diboride is significantly
affected by the microstructural bias, including phase distribution and morphology,
formed during synthesis and processing.
D. In general, the average resistivity of the MOR bars that were cut from the hot
pressed manually mixed powders was lower than the average resistivity of the
MOR bars that were cut from the hot pressed SHS powders.
E. No significant affect on the resistivity could be discerned between the
composites with TiB2 segregated at the alumina grain boundaries (0.213-0.336
ohm-cm) and the composites with TiB2 uniformly distributed amongst the
alumina grains (0.234-0.266 ohm-cm).
F. After the four-hour hold at temperature, the SHS composite resistivity was
comparable to that of the MM composites.
G. Electrical resistivity measurements have the potential of being a non-
destructive means of screening candidate armor materials.
ACKNOWLEDGEMENTS
The author gratefully acknowledges support from the U. S. Army Research
Office Contract No. DAAG55-98-1-0454; Mr. Matthew Burkins at the U. S.
Army Research Laboratory for the ballistic test results; and the U. S. Army
TACOM, Warren MI, Contract No.DAAE07-95-C-R040.
REFERENCES
1 K. V. Logan, “Composite Ceramics,” Final Technical Report, USATACOM,
Warren, MI, Contract #DAAEO7-95-C-R040, November 1996. 2 Ibid 1.
3 K. V. Logan, “Process for Controlling the Microstructural Bias of Multi-Phase
Composites,” U. S. Pat. No 6,090,321, July 18, 2000. 4 K. V. Logan, “Process for Controlling the Microstructural Bias of Multi-Phase
Composites,” U. S. Pat. Notice of Allowability, Application No. 09/549,648,
March 12, 2002. 5 A. R. Keller and M. Zhou, “Effect of Microstructure on Dynamic Failure
Resistance of Titanium Diboride/Alumina Ceramics,” Journal of the American
Ceramic Society, to be published in 2002. 6 G. Kennedy, L. Ferranti, R. Russell, M. Zhou and N. Thadhani, “Dynamic High
Strain-Rate Mechanical Behavior of Microstructurally-Based Two-Phase
TiB2+Al2O3,” Journal of the American Ceramic Society, to be published in 2002. 7 Louis Ferranti Jr., “Processing and Characterization of Microstructurally Biased
Two-Phase Titanium Diboride/Alumina Ceramic (TiB2+Al2O3),” Masters
620 Ceramic Armor Materials by Design
Thesis, School of Materials Science and Engineering, Georgia Institute of
Technology, Atlanta, Georgia, December 2001. 8 L.H. Van Vlack, pg 493 in Elements of Materials Science and Engineering,
Addison-Wesley Publishing Company, Reading, MA, 1985. 9 E. C. Skaar and W. J. Croft, "Thermal Expansion of TiB2," Journal of the
American Ceramic Society, 56 pg 45 [1] (1973). 10
K. V. Logan, G. R. Villalobos, and J. T. Sparrow, "Synthesis/Densification
Using SHS of Composite TiB2/Al2O3," presented at The First International
Ceramic Science & Technology Congress, Anaheim, California, 31 October - 3
November, 1989. 11
Z. Rosenberg, S. N. Brar, et al., "Shear Strength of Titanium Diboride Under
Shock Loading Measured By Transverse Manganin Gauges," presented at The APS
1991 Topical Conference on Shock Compression of Condensed Matter,
Williamsburg, VA, June 17-20, 1991, Elsevier. 12
D. P. Dandekar, "Effect of Shock Reshock on Spallation of Titanium Diboride,"
presented at the APS Topical Conference on Shock Compression of Condensed
Matter, Williamsburg, VA, June 17-20, Elsevier. 13
J. Lankford, "Compressive Strength and Microplasticity in Polycrystalline
Alumina," Journal of Materials Science, 12 791-796 (1977). 14
M. G. Stevenson and P. L. B. Oxley, "Experimental Investigation of the Influence
of Speed and Scale on the Strain-Rate in a Zone of Intense Plastic Deformation,"
Proc. Inst. Mech. Engr., 184, [31] 561-74 (1969-70). 15
W. P. Holbrook, ed., "Technical Data," Ceramic Source, 7, 269-369, (1991-
1992).16
D. Viechnicki, W. Blumenthal, et al., "Armor Ceramics - 1987," Proceedings of
the Third TACOM Coordinating Conference, Monterey, CA (1987). 17
D. P. Dandekar and P. J. Gaeta, "Extent of Damage Induced in Titanium
Diboride Under Shock Loading," pp.1059-1068 in Shock Waves and High Strain-
Rate Phenomena in Materials, Marcel Dekker, NY (1992). 18
K. P. Ananthapadmanbhan, P.V. Sreekumar, “Electrical, Resistivity of Plasma-
Sprayed Titanium Diboride Coatings,” Journal of Materials Science 28, [6],
1665-1658 (March 1993) 19
A. J. Moulson and J. M. Herbert, Electroceramics: Materials, Properties,
Applications. Chapman and Hall, 1990 20
Ibid 1,3,4 21
J. K. Phillips, K. V. Logan and R. Gerhardt, “Effects of Hot Press Parameters
on Microstructure and the Effects of Microstructure on Electronic Properties of a
70% Al2O3/30% TiB2 Composite,” Independent Research Report, Mate 4951-2-3,
Georgia Institute of Technology, 1996.
Ceramic Armor Materials by Design 621
ELECTRICAL RESISTIVITY
0
10
20
30
40
50
60
70
80
M M
500
30
1
SHS
500
30
2
M M
500
150
3
SHS
500
150
4
M M
3375
30
5
SHS
3375
30
6
M M
3375
90
7
SHS
3375
90
8
M M
3375
150
9
SHS
3375
150
10
M M
5/5*
30
11
SHS
5/5*
30
12
M M
5/5*
150
13
SHS
5/5*
150
14
M M
5000
30
15
SHS
5000
30
16
M M
5000
150
17
SHS
5000
150
18
M M
5000
240
21
SHS
5000
240
22
M M
5/5*
240
24
SHS
5/5*
240
23
PRESSURE
HOLD TIM E
SAM PLE NO.
Resistivity (ohm-
0
20
40
60
80
100
120
Density (% theo
RESIST. 1 stat. avg. RESIST. 2 stat. avg. BAR DEN. % theo.
5/5*=500/5000
i
Figure 16. Summary of electrical resistivity measurements
622 Ceramic Armor Materials by Design
PHASE EQUILIBRIUM STUDIES IN AL2O3-TIB2
Isabel K. Lloyd Kevin J. Doherty and Gary A. Gilde
Materials and Nuclear Engineering U.S. Army Materials Research Laboratory
University of Maryland Aberdeen Proving Ground
College Park, MD 20742-2115 Aberdeen, MD 21005
ABSTRACT
In this study, high temperature anneals were preformed on Al2O3-TiB2
mixtures containing 10, 20 and 40 mole % TiB2 to determine if the eutectic
reaction suggested by the microstructure of self-propagating high temperature
synthesis powders and thermodynamic calculations occurred. Energy dispersive
spectroscopy and X-ray diffraction of the annealed mixtures suggested that there
was a eutectic near 1925°C at a composition near 80 mole % Al2O3 and 20 mole
% TiB2. The melting behavior of the mixtures and the microstructures of the
annealed powder mixtures supported this conclusion.
INTRODUCTION
Al2O3-TiB2 has received some attention as a potential ceramic armor material
since it was expected to retain some of the hardness and stiffness of TiB2 while
being easier to process than monolithic TiB2. Interest in the system was piqued by
initial ballistic tests that suggested it exhibited significant resistance to high strain
rate penetration and mechanical properties tests that indicated its static mechanical
properties were similar to TiB2 [1]. Al2O3-TiB2 bodies can be made by hot-
pressing either mechanically mixed Al2O3 and TiB2 powders or composite
powders made by SHS, self-propagating high temperature synthesis, around
1600°C. The microstructure [2,3] of the SHS powders suggests that there may be
a eutectic between Al2O3 and TiB2. This conclusion is supported by the
microstructure of hot-pressed samples that experienced significant temperature
overshoots. This study explored the existence of a eutectic since an understanding
of the phase equilibrium in a system can aid in the development of processing
routes that produce tailored microstructures or that are amenable to large scale
manufacturing.
Before any experiments were done, the potential eutectic composition and
temperature were estimated from freezing point depression and liquidus surface
Ceramic Armor Materials by Design 623
To the extent authorized under the laws of the United States of America, all copyright interests in this publication are the propertyof The American Ceramic Society. Any duplication, reproduction, or republication of this publication or any part thereof, withoutthe express written consent of The American Ceramic Society or fee paid to the Copyright Clearance Center, is prohibited.
calculations. It was assumed that the liquid in the Al2O3-TiB2 system was an ideal
solution. Then, mechanically mixed powders were annealed in W foil packets in
Ar at temperatures between 1850 and 2070°C. After annealing the powder
mixtures were examined visually, optically, in the SEM with EDS (energy
dispersive spectroscopy) and back scattered electrons, and by XRD (X-ray
diffraction).
THERMODYNAMIC CALCULATIONS
Freezing point depression was used to estimate the eutectic composition
and temperature. First the freezing point depression for alumina as a function of
TiB2 addition was calculated assuming an ideal solution using equation 1 [4]:
ln XA=-( Hf/R) [(Tm-T)/( Tm * T)] (1)
where XA is the mole fraction TiB2, Tm is the melting point of alumina (2327 K
[5]), Hf is the enthalpy of fusion (-1675.7 kJ/mol [5]) and R is the Universal Gas
Constant. Then, the freezing point depression of TiB2 was similarly calculated.
Both liquidus curves were plotted as shown in Figure 1 and the temperature and
composition where the two liquidus curves crossed was taken as the estimate of
the eutectic composition and temperature.
The intersection of the liquidus surfaces for both Al2O3 and TiB2 was used
to confirm the estimate from freezing point depression. To calculate the liquidus
surfaces, the change in free energy for ideal solution as a function of temperature
from 2200 to 3100 K was first calculated using equation 2 [4]:
G = RT [X lnX + (1-X) ln (1-X)] (2)
where R is the gas constant, T is the temperature, and X is the mole fraction of
solute. Next, the difference in free energy between the solid and liquid was
calculated using equation 3 [4]:
Gs – Gl = - Hf ln (Tm /T) (3)
where Gs is the free energy of the solid, Gl is the free energy of the liquid and Hf
is the enthalpy of fusion for either Al2O3 or TiB2. A tangent to the G curve
through the Gs – Gl value was then drawn to estimate the liquidus composition at
that temperature. These values are also shown in Figure 1. The intersection of
both sets of curves occurs at about 80 mole % alumina and about 2250 K
(1977°C).
.
624 Ceramic Armor Materials by Design
1800
2000
2200
2400
2600
2800
3000
3200
3400
0 0.2 0.4 0.6 0.8 1
X alumina
T (
K)
Al2O3 (FP dpress)
TiB2 (FP dpress)
Al2O3 (liquidus)
TiB2 (liquidus)
Figure 1: Estimate of eutectic temperature and composition
EXPERIMENTAL PROCEDURE
Mixtures of commercial Al2O3(Alcoa, A16) and TiB2 (Stark, Grade D)
powders with the compositions given in Table I were ball milled in ethanol for 20
hours and then dried under a heat lamp. After drying, large agglomerates were
crushed with a spatula and 2-5 g of mixed powder was placed in a loosely sealed
W foil packet. Then, a packet of each composition was placed in its own covered
graphite crucible and all three compositions were annealed in a graphite furnace
under the conditions in Table II.
After annealing the samples were examined visually and in a optical
microscope for signs of melting and reaction with the W foil. Then they were
examined using XRD with Cu K radiation to determine phase composition. The
strong alumina peaks (2 = 35.2, 25.6, 43.4, 66.5 and 68.3°) in the as-milled,
unannealed powders were used as standards. Quantitative comparison with these
peaks was used to determine the relative amounts of Al2O3 and TiB2 in the
annealed powders. The annealed powder was lightly coated with gold to prevent
charging before it was examined in the scanning electron microscope (SEM) using
secondary and back-scattered electrons and EDS.
Ceramic Armor Materials by Design 625
Table I: Compositions used for eutectic studies
Al2O3 TiB2 Predicted Melting Behavior
mole% wt.% mole% wt/%
80 (85.4) 20 (14.6) predicted eutectic composition
90 (93) 10 (7) 50 m% eutectic L, 50 m% Al2O3
60 (70) 40 (30) 50 m% eutectic L, 50 m% TiB2
Table II: Heat Treatments
Run Temperature (°C) Time Atmosphere
1 1850 4 hours vacuum
2 1900 4 hours Ar
3 2070 15 min. Ar
4 1950 4 hours Ar
5 1925 4 hours Ar
6 1925 4 hours Ar
RESULTS AND DISCUSSION
Evidence of melting was observed visually, under the optical microscope,
and in the SEM for samples annealed at 1925°C and above. No melting was
observed below 1925°C. More melting was observed in the 80 mole % alumina
mixtures. The microstructures were not uniform in the samples annealed
1925°C. In these samples, EDS indicated that the top of the powder bed was
highly deficient in alumina and that the bottom of the powder bed was less
alumina deficient. The 80 mole % alumina samples were more alumina deficient
than the other samples. These observations were supported by the quantitative
XRD results shown in Figure 2. Since alumina would be more volatile in the
liquid state than the solid state, this was taken as evidence that there was more
alumina in the liquid state in the 80 mole % mixtures. XRD indicated that Al2O3
and TiB2 were the major phases in all samples. Later experiments used an Ar
atmosphere since a number of minor phases were formed under vacuum in the
first experiment. There was no evidence of reaction with the W foil in the XRD
results.
The microstructures observed for both polished and as-annealed powders
were consistent with the proposed eutectic. Below 1925°C, the TiB2 grains were
angular and separated by an alumina matrix. Above 1925°C, as shown in Figure
3, the TiB2 grains were rounded and they tended to be more interconnected which
indicates melting and suggests the possibility of a eutectic.
626 Ceramic Armor Materials by Design
0
10
20
30
40
50
60
70
80
90
100
1700 1800 1900 2000 2100
Anne aling T (°C, as re c 'd =1800)
80mole%alumina
(peak at 68.3°)80mole%alumina
(peak at 43.4°)80mole%alumina
(peak at 35.2°)80mole%alumina
(peak at 25.6°)80mole%alumina
(peak at 66.5°)
Figure 2: Semi-quantitative XRD results for 80 mole% alumina. Ratios were
calculated using different alumina peaks in the as-milled powder as standards.
Figure 3: Secondary electron image of an 80 mole % alumina powder
mixture annealed at 1950°C. The rounded dark grains are TiB2.
Ceramic Armor Materials by Design 627
CONCLUSIONS
Thermodynamic calculations and annealing experiments in the Al2O3-TiB2
system from 60 to 90 mole % Al2O3 indicate a liquidus minimum above 1925°C
and a possible eutectic at about 80 mole % Al2O3. Additional experiments will be
needed to confirm these tentative conclusions. Given the high liquidus
temperature of the proposed eutectic, it is unlikely that eutectic processing would
offer any advantages for Al2O3-TiB2 composites compared to conventional hot-
pressing [1].
REFERENCES
1. K. V. Logan, “Composite Ceramics,” Final Technical Report A002, Army
Materials and Mechanics Research Center Contract DAAE07-95, Nov.
1996.
2. K.V. Logan and J.D. Walton, “Ti Formation Using Thermite Ignition,”
Ceram. Eng. Proc. 5 [7] 712-38 (1985).
3. L.J. Kecskes, A. Niiler, T. Kottke, K.V. Logan, and G.R. Villalobos,
“Dynamic Consolidation of Combustion Synthesized Alumina-Titanium
Diboride Composite Ceramics,” J. Am. Ceram. Soc. 79 [10] 2687-95
(1996).
4. C.F. Bergeron and S.H. Risbud, Phase Equilibrium Studies in Ceramics,
pp. 52-62, Am. Ceram. Soc. 1984.
5. JANAF Tables Al2O3, TiB2, J. Phys. Chem. Ref. Data, Monograph 9
628 Ceramic Armor Materials by Design
MICROSTRUCTURE DEVELOPMENT OF ALUMINUM OXIDE/TITANIUM
DIBORIDE COMPOSITES FOR PENETRATION RESISTANCE
J.W. Adams, G.A. Gilde and M. Burkins
U.S. Army Research Laboratory
Aberdeen Proving Ground, MD 21005
L. Prokurat Franks
U.S. Army Tank-Automotive and Armaments Command
Warren, MI 48397
ABSTRACT
Early research on aluminum oxide/titanium diboride (Al2O3/TiB2) composites
focused on exploiting their potential as a low cost armor ceramic. Limited
ballistic data demonstrated that the microstructure has a dramatic effect on
ballistic performance. With the "preferred" microstructure, the penetration
resistance of Al2O3/TiB2 approached that of monolithic TiB2 ceramics. Challenges
were encountered both in quantifying the microstructural detail and fabricating
the desired microstructure.
Our research focused on microstructure control during fabrication and
correlation of microstructure with mechanical properties and penetration
resistance of the composite. Composites were made from mixed Al2O3 and TiB2
powders, as well as a composite Al2O3/TiB2 powder prepared via a self-
propagating high-temperature synthesis (SHS) reaction. A summary of depth of
penetration ballistic analyses for several projectiles is given. Our results show that
although the penetration resistance of Al2O3/TiB2 composites is good, the results
fall within the expected experimental scatter shown by commercial state-of-the-
art armor ceramics.
BACKGROUND
A brief history of the interest in Al2O3/TiB2 composite materials shows that in
1982 the Army became aware of Soviet technology for Self-Propagating High
Temperature Synthesis (SHS) to produce TiB2. Within ten years fully dense SHS
Al2O3/TiB2 composites were produced at Georgia Tech.1 Ballistic evaluations of
those materials performed by the University of Dayton Research Institute showed
that there could be composite a possible correlation between microstructure and
Ceramic Armor Materials by Design 629
To the extent authorized under the laws of the United States of America, all copyright interests in this publication are the propertyof The American Ceramic Society. Any duplication, reproduction, or republication of this publication or any part thereof, withoutthe express written consent of The American Ceramic Society or fee paid to the Copyright Clearance Center, is prohibited.
ballistic performance. The composite structure that had TiB2 localized at the grain
boundaries of the aluminum oxide exhibited a high mass efficiency. In 1997
TARDEC contracted with the Army Research Laboratory for further study based
on the following factors:
•Promising new materials processing technology
•Ballistic test results showed potential
•Comparable performance to state-of-the-art armor ceramics, with
possible cost savings
•Potential Future Combat System (FCS) applications
In particular, the early Al2O3/TiB2 composites that were processed using
powders derived by self-propagating synthesis (SHS) and evaluated by long-rod
penetrators in Depth of Penetration (DOP) tests gave intriguing results. The
ballistic mass efficiencies were greater than expected from the rule of mixtures,
and were high enough to generate interest in these materials as potential armor
(see Figure 1.) 2,3
The reason for interest in the composite is twofold: 1) initial
screening indicated that the material may perform as well as, or equivalent to
titanium diboride (TiB2) armor ceramics at substantially lower cost, and
630 Ceramic Armor Materials by Design
Figure 1. Early DOP ballistic data for L/D=13 rod at 1550 m s-1
against alumina-
titanium diboride ceramics.
2) Al2O3/TiB2 has higher space efficiency than can be achieved with silicon
carbide, as well as a mass efficiency that is almost equal to that of silicon carbide
against medium caliber threats.
The purpose of our work had several aspects. We wanted to explore other
powder processing and sintering routes to systematically determine and quantify
differences in microstructures, to evaluate the composites against small and
medium caliber penetrators and assess consistency with previous ballistic results,
and lastly, to correlate the microstructure to the ballistic properties.
EXPERIMENTAL
We fabricated composites according to several processes, characterized them
and performed DOP ballistic tests using three different penetrators in the course
of investigating this system. Details of the processing matrix, mechanical and
DOP evaluation methods and test results are given in Gilde et al.4 Processes used
in this study were:
• SHS Al2O3/TiB2 powders + ball milling + hot pressing (HP)
• Al2O3 and TiB2 powders + ball milling + HP
• Co-extrusion Al2O3 and TiB2 powders + HP
• Colloidal powders + sintering
• Al2O3 and TiB2 powders shaken to mix electrostatically (ESD) + HP
Our processing study maintained the nominal 75/25% composite ratio weight
ratio of Al2O3/TiB2 using various green powder processing routes to achieve
different microstructural textures. Density, hardness, 4-point flexure strength,
fracture toughness and fracture analyses were performed on all composites prior
to ballistic evaluation. DOP ballistic testing was conducted using 7.62 mm AP M2
(armor piercing) projectiles and L/D=10 tungsten alloy rods. However, the larger
L/D=13 tungsten alloy rods that had been used in the 1990s were not a part of this
study.
RESULTS
Figure 2 summarizes the results of the ballistic testing against the 7.62mm AP
M2 projectile and compares it to the ballistic performance of hot pressed silicon
carbide, boron carbide and a sintered aluminum oxide tested against the same
penetrator.5 As can be seen from the graph, the aluminum oxide/titanium diboride
composites performed slightly better than the sintered aluminum oxide, but were
less effective than silicon carbide and boron carbide. The composites made from
the SHS powders performed the best against this small projectile.
Ceramic Armor Materials by Design 631
Penetr
ation
resis
tan
ce x
Are
aldensity, kg m
-2
0
20
40
60
80
100
120
140
0 5 10 15 20 25
Areal density, kg m-2
Al2
O3
: 123.12 + 13.590(1-e (0.11659)C
tC
) R2 = 0.9864
B4
C: 123.12 + 5.1494(1-e (0.3086)C
tC
) R2 = 0.98066
SiC: 123.12 + 25.113(1-e (0.13106)
C
t
C
) R2 = 0.97985
FM CP
SHSMM CP
Al2O3/TiB2 composites
Figure 2. Residual penetration areal density vs. ceramic areal density against the
7.62 AP M2 projectile at 841 m s-1
.
Ballistic properties of Al2O3/TiB2 composites impacted with a 131W tungsten
alloy rod at 1550 m/s as compared to other armor ceramics are presented in Table
2. Several preferred Al2O3/TiB2 microstructures were evaluated. Despite the
overlap in the ballistic data for aluminum oxide, titanium diboride and silicon
carbide, the average em follow the expected trend that silicon carbide is better than
titanium diboride, and titanium diboride performs better than aluminum oxide. It
Table 2. Ballistic properties of armor ceramics impacted with a L/D=13 rod at
1550 m s-1.
Sample em es q2
Commercial AD995 3.2 + 0.5 1.6 + 0.3 5.0 + 1.9
Commercial SiC 4.4 + 0.5 1.8 + 0.3 8.0 + 2.3
Commercial TiB2 3.9 + 0.3 2.2 + 0.2 8.7 + 1.6
SHS TiB2 around Al2O3 3.3 + 0.2 1.7 + 0.1 5.6 + 0.6
MM TiB2 around Al2O3 4.1 + 0.4 2.1 + 0.3 8.7 + 2.1
SHS TiB2 within Al2O3 2.5 + 0.1 1.3 + 0.1 3.4 + 0.2
MM TiB2 within Al2O3 3.3 + 0.6 1.7 + 0.4 5.5 + 2.7
632 Ceramic Armor Materials by Design
can be seen that the manually mixed (MM) Al2O3/TiB2 composites have higher em
values than AD 995 aluminum oxide. The average em is higher than that of
titanium diboride and slightly less than hot-pressed silicon carbide. When both
space and weight are critical to the armor design, the es and q2 values indicate that
the Al2O3/TiB2 composite could have an advantage over titanium diboride armor
ceramics for an armor package designed against medium cal threats.
Ballistic properties of Al2O3/TiB2 composites impacted with the L/D=10
tungsten alloy rod at 1500 m s-1
as compared to other armor ceramics are
presented in Figure 2 and Table 3.
x
0
5
10
15
20
25
30
35
40
45
40 60 80 100 120 140 160 180
Ceramic areal density, kg m-2
Pe
ne
tra
tio
n,
mm
HP SiC 25mm
HP SiC 30mm
TARDEC GTRI
TARDEC ARL ESD
1992 GTRI UDRI 39MM
X HP TiB2
SiC
Al2O3/TiB2
xx
x
Figure 2. Penetration into RHA backing vs. ceramic areal density (25 mm
thickness) against the L/D=10 rod at 1500 m s-1.
Table 3. Ballistic properties of armor ceramics (25 mm thickness) impacted with
a L/D=10 rod at 1500 m s-1.
Material Density, gcm-3
em
Commercial AD995 3.6 2.4
Commercial SiC 3.2 4.2
Commercial TiB2 4.5 3.2
Al2O3/TiB2 (ESD) 4.1 3.1
All Al2O3/TiB2 composites 4.1 3.0
Ceramic Armor Materials by Design 633
SUMMARY
Our investigation to assess Al2O3/TiB2 composites’ potential as an armor
ceramic demonstrated that distinctive microstructural textures can be developed
and controlled by a variety of processing methods. A systematic ballistic
evaluation was completed for small and medium caliber projectiles at velocities
ranging from ~850 m/s to 1500 m s-1
. All TiB2/ Al2O3 composite structures were
effective at defeating the projectile in all cases. For the 7.62 AP round, the
composites made from the SHS powder performed slightly better. In the case of
medium caliber long rod penetrators, SHS-derived composites did not offer any
advantage. Composites made from mixed Al2O3 and TiB2 powders performed
better. The process of mixing dry powders to electrostatically disperse the TiB2
around the Al2O3 grains resulted in composite structures that were as effective as
those that were ball milled for hours.
In order for a ceramic to offer attractive potential as armor, the material must
offer effective protection, and be manufacturable and affordableThe promise
based on early ballistic data and probable cost savings for Al2O3/TiB2 composites
has not been borne out in this study. Serious manufacturability issues, including
the lack of commercial SHS powder suppliers and little market pull for products
beyond armor for the titanium diboride plus alumina system, override the
estimates for favorable raw material cost/processing savings. Early anecdotal high
ballistic penetration resistance results were shown to be within the range of
expected DOP test variability.
REFERENCES 1 K.V. Logan, “Elastic-plastic Behavior of Hot-pressed Composite Titanium
Diboride/Alumina Powders Produced Using Self-propagating High-temperature
Synthesis,” PhD Thesis, Georgia Institute of Technology, 1992. 2 G. Abfalter, N.S Brar. and D. Jurick, “Determination of the Dynamic
Unload/Reload Characteristics of Ceramics,” University of Dayton Research
Institute, Dayton OH, June 1992, Contract No. DAAL03-88-K-0203. 3 P. Woolsey, D. Kokidko and S. Mariano, “An Alternative Test Methodology
for Ballistic Performance Ranking of Armor Ceramics,” MTL TR 89-43, U.S.
Army Materials Technology Laboratory, Watertown, MA, 1989. 4G.A. Gilde, J.W. Adams, M. Burkins, M. Motyka, P.J. Patel, E. Chin, L.
Prokurat Franks, M.P. Sutaria and M. Rigali, "Processing of Aluminum
Oxide/Titanium Diboride Composites for Penetration Resistance," Cer. Eng. Sci.
Proc., 22 (2001) 331-342.5 T. J. Moynihan, S. Chou, and A.L. Mihalcin, “Application of the Depth-of-
Penetration Test Methodology to Characterize Ceramics for Personnel Protection”
ARL-TR-2219, April 2000, Army Research Laboratory, Aberdeen Proving
Ground, MD 21005-5066.
634 Ceramic Armor Materials by Design
THE EFFECT OF METAL-CERAMIC BONDING ON BALLISTIC IMPACT
Kevin J. Doherty
US Army Research Laboratory
Weapons Materials Research Directorate
AMSRL-WM-MC
Aberdeen Proving Ground, MD 21005
ABSTRACT
Lightweight armor systems are crucial to the survivability of future Army
vehicles. The combination of ceramics and lightweight metals is a key element in
modern armor packages. The interface created when joining metals and ceramics
can have a significant influence on the behavior of the entire system. In this
study, the joining of SiC and Ti-6Al-4V plates was demonstrated using an active
solder, Sn-4Ag, containing ~4 wt% Ti. This configuration was compared with
plates joined using an epoxy. Preliminary ballistic evaluation and microstructural
analysis of the joints in the different armor systems will be discussed.
INTRODUCTION
The desire for smaller, lighter Army vehicles has motivated the need for
lightweight metal and ceramic armor systems. The process of fabricating an
armor package from lightweight metals and ceramics is complicated by the need
to bond very dissimilar materials both together as well as attaching these armor
packages to the vehicle structure. A typical joining method for ceramics-metals is
adhesive bonding. Joining with adhesives, such as epoxy, is convenient because it
is performed near room temperature, in air and is compatible with most materials.
The drawbacks to adhesive bonding are the resulting low bonding strength and the
low modulus. The combination of low modulus and low density creates a
substantial elastic impedance mismatch with the ceramic and metal substrates.
Other bonding options such as brazing and soldering typically have higher moduli
and higher densities that decrease the elastic impedance mismatch with the
ceramic and metal substrates in comparison with adhesives.
The desire for stronger bonding in metal-ceramic systems has led to the
examination of joining techniques that involve beneficial chemical reactions at the
Ceramic Armor Materials by Design 635
To the extent authorized under the laws of the United States of America, all copyright interests in this publication are the propertyof The American Ceramic Society. Any duplication, reproduction, or republication of this publication or any part thereof, withoutthe express written consent of The American Ceramic Society or fee paid to the Copyright Clearance Center, is prohibited.
metal-ceramic interfaces. During the early ‘80s, Mizuhara and coworkers [1,2]
adapted this idea from the ‘50s by putting an “active” component, such as
titanium (Ti), directly into a brazing alloy, typically a silver-copper eutectic, to
significantly improve the wetting of both metal and ceramic substrates. This
initiates a one-step vacuum brazing process that wets most materials (including
ceramics, Ti alloys and stainless steels) and forms strong, metallurgical bonds.
The major disadvantage in using “active” brazing for metals and ceramics is the
high processing temperature required that results in large strain (stress) build-up
from the inherent differences in coefficient of thermal expansion (CTE) between
metals and ceramics during cooling. There are some techniques available to
alleviate the strains on the ceramic, such as using an interlayer, which either has
an intermediate (between the metal and ceramic) value of CTE and/or is “soft”
(compliant). However, it is still extremely challenging to actively braze specimens
that are larger than 5 cm in diameter when there is a considerable CTE gradient.
Active solder joining is an emerging technology that incorporates many of
the ideas from active brazing. A reactive element (typically Ti) is added to a
solder alloy to enable direct wetting and bonding. Currently, two lead-free
systems are being investigated: Sn-Ag-Ti and Zn-Ag-Ti [3]. No chemical fluxes
are used, so mechanical agitation (such as brushing or ultrasonic vibration) is used
to disrupt the oxide naturally on the solder to promote wetting. The use of lead-
free solders, without chemical fluxes, creates an “environmentally friendly”
process that offers additional cost benefits by eliminating extra cleaning steps
associated with the fluxes. The active soldering process offers a compromise of
lower joining temperatures (<450ºC), reasonable elastic impedance, but little
improvement in strength over epoxy. The flexibility of this process allows the
joining of larger specimens with a significant mismatch in CTE, which offers the
chance to test the effect of elastic impedance and adherence of the bonding layer
on ballistic samples. Thus, this paper will examine the effect of active soldering
and epoxy joining for SiC and Ti-6Al-4V substrates and it will present some
preliminary ballistic results for the different bonding conditions.
EXPERIMENTAL
Square ballistic targets were assembled by joining hot pressed silicon
carbide (10.2 cm x 10.2 cm) to annealed Ti-6Al-4V (MIL-DTL-46077F, 15.2 cm
x 15.2 cm). The elastic and bulk moduli of the SiC plates were measured by the
pulsed excitation method (ASTM C1259) to verify the homogeneity within the
lot. The density of each plate was also determined to eliminate any anomalous
plates from the study. The SiC plates were used in an as-ground condition and the
Ti-6Al-4V pieces were machined to achieve flat and parallel surfaces. Four
different bonding conditions were investigated to join the SiC to the Ti-6Al-4V.
636 Ceramic Armor Materials by Design
The first set of targets was bonded using a two-part epoxy (Epon resin 815 and
Versamid 125). The epoxy was applied to the mating surfaces and the targets
were cured for 72 hours at room temperature. The other three sets of targets were
prepared and bonded using the proprietary active solder S-Bond™ Alloy 220 from
Materials Resources International in Lansdale, PA. The S-Bond™ Alloy 220 is a
Sn-4Ag based alloy with 4 weight percent Ti. The first set of SiC and Ti-6Al-4V
plates were grit blasted prior to joining in air with the active solder at 250ºC. The
next set of SiC plates was pre-treated with the active solder in a vacuum furnace
at 850ºC for one-half hour and furnace cooled. The Ti-6Al-4V was grit blasted
and then bonded to the pre-treated SiC in air at 250ºC. The final set of SiC and
Ti-6Al-4V plates were all pre-treated in a vacuum furnace at 850ºC and furnace
cooled prior to final joining in air at 250ºC. An example of the ballistic targets is
shown in Figure 1.
Figure 1. Sample ballistic targets
Projectiles were fired at various velocities at the targets to establish a
modified V50 protection ballistic limit (MIL-STD-662E, V50 Ballistic Test For
Armor). Modifications from the standard procedure included using x-ray analysis
instead of a witness plate for the determination of partial penetration (PP) versus
complete penetration (CP). The targets bonded with epoxy were tested first to
establish the baseline V50 for comparison with the various soldered targets. While
the epoxy targets provided an initial testing velocity, there was still an insufficient
number of pre-treated solder targets to achieve a statistically significant V50. A
photograph of the ballistic setup is shown in Figure 2. The V50 ballistic limit may
be defined as the velocity at which 50% of the attacking projectiles may be
statistically expected to completely penetrate the target.
Optical metallography was performed on the impacted Ti-6Al-4V plates
and SiC fragments. In addition, scanning electron microscopy (with a Robinson
backscatter detector) was utilized to investigate the impacted interface structure.
Hardness measurements were taken from segments of the Ti-6Al-4V plates
Ceramic Armor Materials by Design 637
following the different bonding procedures to determine any consequence from
the heating and cooling cycles.
Figure 2. Sample ballistic setup
RESULTS
The ballistic testing produced comparable V50 results for all of the
different bonding conditions. In all cases, the ballistic event produced extremely
fragmented SiC pieces detached from the Ti-6Al-4V plates. The Ti-6Al-4V plates
can be separated into two categories of ballistic damage. The first set of plates
has undergone PP (Figure 3) where the projectile did not penetrate the target. A
plastically deformed bulge in the center that is embedded with SiC rubble
characterizes these plates. In addition, the plates contain several cracks, but are
still intact. The second set was CP (Figure 4), characterized by slight bulge in the
center with a crack proceeding entirely through the center of the plates, from top
to bottom. All of the Ti-6Al-4V plates contain scratches and an adhered powdery
residue emanating from the center out to the edges.
SiC fragments were collected and characterized after the ballistic tests.
Typically, the fragments examined came from the corners and edges of the
ballistic targets. Optical microscopy was utilized to examine the SiC/solder
interface in the grit blasted and soldered condition. Only a small fraction of the
solder is still adhered to the SiC following the ballistic event (indicated by the
lighter regions in Figure 5). Optical images of the SiC fragments, joined by
soldering with a pre-vacuum treatment, are presented in Figure 6. For this joining
condition the solder is still adhered to the surface of the SiC. The raised regions
(islands) in Figure 6b constitute failure at or near the Ti-6Al-4V/solder interface,
while the recessed regions constitute failure within the solder. This behavior is
638 Ceramic Armor Materials by Design
further demonstrated in a backscatter SEM image (Figure 7) of another SiC
fragment from the same ballistic event as Figure 6. The bonding of the solder to
the SiC is fully intact while the failure is both within the solder and at or near the
Ti-6Al-4V/solder interface.
Figure 3. Ti-6Al-4V plates after PP. Bonded with (a) epoxy, (b) solder, and (c)
solder with vacuum treatment.
Figure 4. Ti-6Al-4V plates after CP. Bonded with (a) epoxy, (b) solder, and (c)
solder with vacuum treatment.
Figure 5. Optical images of post-ballistic SiC, grit blasted and bonded with solder.
Ceramic Armor Materials by Design 639
Hardness measurements taken from the Ti-6Al-4V plates after the epoxy
or grit blast/solder bonding produced results consistent with the standard MIL-
DTL-46077F (33.0-33.5 HRC). With the addition of a vacuum treatment to
promote improved bonding, the softening of the Ti-6Al-4V to 32 HRC was
measured and attributed to the slow furnace cool.
Figure 6. Optical images of post-ballistic SiC, soldered with vacuum treatment.
Figure 7. SEM backscatter image, post-ballistic SiC soldered, vacuum treatment.
DISCUSSION
Using active soldering offers a new alternative for joining large samples of
materials where there is a significant CTE mismatch, such as SiC and Ti-6Al-4V.
The obvious contrasts between epoxy and the Sn-Ag-Ti alloy offers an excellent
opportunity to explore both material property differences (Table I) and bonding
differences as they relate to ballistics. The increased elastic modulus and density
of the active solder results in a tenfold increase in the elastic impedance over the
epoxy. This increase brings the elastic impedance more in line with the SiC,
decreasing the impedance mismatch, and causing less reflection of the stress wave
back into the SiC during a ballistic event. The different processing routes used for
active soldering of the SiC and Ti-6Al-4V also allows for a variation in bond
strength. Active soldering without a thermal treatment provided the lowest bond
640 Ceramic Armor Materials by Design
strength of all the techniques. In the epoxy and the grit blast/solder cases, the
bond strength comes exclusively from van der Waals forces, which is verified in
the post-impacted images. The thermal treatment improves adherence of the
solder to the SiC resulting from chemical reactions between the titanium in the
solder and the SiC. Thermal treatment of the substrates prior to active soldering
can improve the bond strength by 40% [4]. However, even with the variations in
elastic impedance and bond strength among the different joining techniques, the
ballistic results show no difference in performance in any of the configurations.
Table I. Typical properties for joining materials and substrates***^^^"
MaterialDensity
(g/cm3)
Elastic
Modulus
(GPa)
Long.
Velocity"
(km/s)
Elastic
Impedance"
(kg/m2s)*10
7
CTE
(m/m°C)
[RT-x°C]
Tensile
Strength
(MPa)
Bonding
Temp.
Epoxy* Epon 815 1.1 3-4 1.7 0.2 73 <100°C
Solder**S-Bond™
Alloy 2207.4 50-56 2.7 2
19*10-6
(250°C)53 250°C
Ceramic^Hot Pressed
SiC3.2 455 11.9 3.8
4.5*10-6
(1000°C)
Metal^^Annealed
Ti-6Al-4V4.4 114 5.1 2.3
9.7*10-6
(650°C)
Grit blasting of the SiC surface is the initial attempt to improve the
adherence of the solder to the SiC. Typically, grit blasting adds surface roughness
that increases bond area and introduces mechanical interlocking. However, in
post-ballistic analysis the adherence is limited and is generally confined to the
depressions in the SiC surface (Figure 5). The thermal treatment of the SiC prior
to active soldering leads to a substantial improvement of the adherence of the
solder to the SiC (Figure 6) which relates to an increase in the shear strength of
the bond. Such an improvement can bring the shear strength of the bond above
the tensile strength of the solder. This is evident in Figures 6 and 7 where the
failure of the bond occurs within the solder and near the solder/Ti-6Al-4V
interface instead of at the SiC/solder interface.
The thermal treatment did not have the same remarkable effect with the
Ti-6Al-4V as it did on the SiC surface. Active soldering of grit blasted Ti-6Al-4V
* Epoxy information from Resolution Performance product data sheets.
** Solder information from Materials Resources International product data sheets.
^ SiC information from Cercom, Inc. product data sheets.
^^ Ti-6Al-4V information from TIMET product data sheets.
" Longitudinal velocity and elastic impedance values calculated from the above data.
Ceramic Armor Materials by Design 641
is still a van der Waals type bond. The thermal treatment of the Ti-6Al-4V did
not induce the level of reaction between the active titanium in the solder with the
Ti-6Al-4V surface that was observed in the SiC. This behavior occurred because
the active titanium most likely reacted more with the protective oxide on the Ti-
6Al-4V surface and reacted less with the base metal.
Even though there was not a correlation between ballistic results and better
elastic impedance or improved bond strength, some basic questions are still
relevant. Were the changes in joint properties too insignificant to effect the
ballistics? Would more substantial changes in strength and/or elastic impedance
correlate with better ballistic performance? The use of higher temperature active
solders and brazes should allow experiments to be performed to answer these
questions. However, some basic material changes may be required in order to
diminish the higher strains (stresses) associated with the higher temperature
joining. In addition, quantification of joint shear strength for all of the different
bonding techniques is required to directly relate the bonding and the ballistics.
SUMMARY
• Active soldering is a viable option for joining large pieces of ceramic and
metal where there is a substantial coefficient of thermal expansion mismatch.
• No relationship was observed between ballistic performance and either
bond adherence or elastic impedance in this set of experiments.
• A vacuum thermal treatment prior to joining improves the solder
adherence to the SiC significantly, but has a small effect on the adherence of the
solder to Ti-6Al-4V.
• Higher temperature active solders and brazes allow a larger variation in
bond strength and elastic impedance to further test the relationship between
bonding and ballistics.
ACKNOWLEDGEMENTS
The author would like to thank Dr. Ernest Chin and Dr. Joseph Wells for
reviewing the manuscript and discussing the research.
REFERENCES 1H. Mizuhara and K. Mally, “Ceramic-to-Metal Joining with Active
Brazing Filler Metal,” Welding Journal, 64 [10] 27-32 (1985). 2H. Mizuhara and E. Heubel, “Joining Ceramic to Metal with Ductile
Active Filler Metal,” Welding Journal, 65 [10] 43-51 (1985). 3R.W. Smith, “Active Solder Joining of Metals, Ceramics and
Composites,” Welding Journal, 80 [10] 30-35 (2001). 4R.W. Smith, Personal Communication, October, 2001.
642 Ceramic Armor Materials by Design
ASPECTS OF GEOMETRY AFFECTING THE BALLISTIC PERFORMANCE
OF CERAMIC TARGETS
I M Pickup, A K Barker, R Chenari and B J James
Defence Science and Technology Laboratories
Chobham Lane, Chertsey
Surrey, KT16 0EE, UK.
V Hohler, K Weber and R Tham
Faunhofer-institut fur Kurzzeitdynamik (EMI)
Eckerstrasse 4,
79104 Freiburg, Germany
ABSTRACT
Some ceramic armour configurations have the ability to erode kinetic energy
long rod penetrators on the impact surface either totally or partially before
subsequent penetration. This phenomenon (often called dwell) can result in very
high ballistic efficiency. The occurrence of dwell is very sensitive to subtle
changes in experimental conditions leading to extreme variation in performance.
The effect of some geometrical parameters such as obliquity, target thickness and
impact surface configuration on ballistic performance of silicon carbide is
assessed at impact velocities ranging from 1400 to 1800 ms-1
. Significant benefits
in ballistic performance may be realised by addressing impact surface and
ceramic back surface configurations to maintain reproducibly high ballistic
performance.
INTRODUCTION
Ceramic armour is capable of exhibiting very high ballistic efficiencies
against kinetic energy (KE) long rod penetrators. Some non-oxide ceramics such
as boron carbide and silicon carbide have a high potential to cause the KE rod to
dwell at the impact surface, i.e. the rod may be eroded at the impact surface
without penetration. This is due to the relatively low density and the particularly
high initial deviatoric strength of ceramics under hydrostatic pressure and in some
cases their relatively slow damage kinetics [1,2]. According to the exact nature of
the experimental targets, KE rods have been totally eroded at the impact surface at
Ceramic Armor Materials by Design 643
To the extent authorized under the laws of the United States of America, all copyright interests in this publication are the propertyof The American Ceramic Society. Any duplication, reproduction, or republication of this publication or any part thereof, withoutthe express written consent of The American Ceramic Society or fee paid to the Copyright Clearance Center, is prohibited.
velocities in excess of 1600 ms-1
[3]. The ballistic efficiency of ceramic armours
can be very significantly increased if such dwell is harnessed reproducibly.
The dwell phenomenon is difficult to quantify. However, there has been some
good experimental measurement of dwell and subsequent KE penetration rates of
slender KE rods into various ceramics by Lundberg et al [4-6] and Orphal and
Franzen [7,8]. This was achieved using multiple flash X-ray photography of
reverse ballistic shots in which the target is fired at the rod. This allowed the
assessment of dwell period as a function of impact velocity for this very specific
experimental situation.
To allow design of improved armour which utilises dwell, it is necessary to
understand what promotes and terminates dwell in a wide range of geometrical
configurations. This precludes the use of high intensity flash X-ray on the grounds
of cost. The work presented here is the initial part of such a study. At this stage
the experimental techniques for determining dwell parameters directly are not yet
mature. Consequently other measurands are employed to evaluate the armour
configuration effects. These are based on a comparison of residual depth of
penetration (DOP) into armour steel back blocks with an estimated penetration in
the absence of dwell.
The ceramic used in this programme is silicon carbide, SiC PAD-B,
manufactured by Cercom Inc., USA. UK specification RHA steel back blocks
were used for DOP measurement. The parameters examined here are:
i) Obliquity. Normal impact and 60 impact angles are compared for
equivalent line of sight (LOS) ceramic thickness.
ii) Velocity. KE rod velocities ranging from 1200 ms-1
to >1800 ms-1
were
employed.
iii) Ceramic thickness. Two LOS thicknesses were used, 30 and 40 mm.
iv) Impact surface configuration. Experiments were conducted with and
without a front coverplate.
EXPERIMENTS
Flat ended tungsten alloy rods (Plansee, Densimet FNC, density = 17600
kgm-3
) with an aluminium flare were fired from a 40 mm smooth bored gun using
a base pushed launch assembly with a three part sabot. The rods were 5 mm
diameter and 100 mm in length. A gun muzzle to target distance of 10 m was used
with 2 pairs of flash X-ray heads positioned 0.1 m and 0.5 m from the target to
monitor rod velocity, pitch and yaw.
All targets were laterally confined using steel adjustable clamps. Annealed
brass inserts were used as an interface between the ceramic and the steel
confinement frame. This was to ensure excellent mechanical contact and
consequently to improve the acoustic impedance match between the confinement
and the ceramic.
644 Ceramic Armor Materials by Design
The normal impact targets and the 60 obliquity targets had lateral dimensions
of 100 x 100 mm and 100 x 200 mm respectively. Front and back surfaces were
ground flat and parallel to 0.01 mm, as were the steel DOP backblocks. The depth
of penetration was assessed by machining the backblocks to determine the
maximum penetration. The cover plate system, where used, was based on a
system used by Hauver et al [3] to accommodate dwell by allowing the possible
lateral spread of rod material as it dwells on the ceramic surface. It consisted of 5
mm of RHA, 2 mm of copper and 1.5 mm of graphite.
RESULTS
Tabulated results are presented in Table I. These identify the experimental
configuration, impact velocity, resolved yaw and measured depth of penetration
for all the shots. The system ballistic mass efficiency, Em, calculated as in
Equation 1, where AD is areal density, is also tabulated.
ceramicDOPRHA
ArefernceRHm
ADAD
ADE (1)
The effect of obliquity on DOP over a range of velocities is presented for 40
and 30 mm thick ceramic targets in Figures I and II respectively. Figures III and
IV show the effect of target thickness for normal and oblique rod impact
respectively. The effect of cover plate addition for both thickness targets is
presented in Figures V and VI.
DISCUSSION
Two methods are employed to provide a reference DOP vs. velocity curve
which represents the DOP after penetration through the ceramic if no penetrator
surface dwell occurred. No direct methods are available so the following
approximations have been made. The first was calculated using steady state
penetration data based on flash X-ray photography [4] from impacts of the same
rod and silicon carbide material and the same velocity regime as in the
experimental programme presented here. Assumptions were made that the same
penetration rates would be applicable to the 5 mm rod, 20:1 aspect ratio (the rods
from reference 4 were 2 mm diameter and 40:1 aspect ratio), and that the tail
velocity was reduced by an amount equal to twice the particle velocity in the rod,
each time an elastic rebound reached the free surface of the tail. Using these
assumptions, tip and tail velocities were calculated as a function of penetration
into the ceramic and the residual length of the rod and its velocity upon entering
the RHA DOP block were estimated. The residual penetration into RHA from this
starting point was calculated by numerical simulation using the Lagrangian code
ELFEN. Johnson-Cook models were used for the rod and for the RHA backblock.
This reference curve is identified in the Figures as ‘no-dwell calculation’.
Ceramic Armor Materials by Design 645
TABLE I. Experimental results
Shot ID Target Type Obliquity Yaw Impact Residual Em
Velocity DOP
( ) ( ) (m/s) (mm)
3094 0 2.1 1593 83.0 0.97
3095 0 2.0 1756 97.0 1.01
3096 0 1.8 1463 69.0 0.97
3098 0 0.6 1200 37.0 1.06
3103 0 2.0 1833 104.0 1.02
3199
RHA
reference shots
0 1.6 1416 61.9 1.00
3204 0 1.6 1630 31.1 1.54
3205 0 1.7 1560 0.0 3.25
3206 0 2.1 1770 0.0 4.19
3315 0 0.2 1708 34.7 1.59
3207 0 - 1717 9.6 2.81
3265 0 2.5 1691 17.1 2.23
3251
40mm SiCB
with coverplate +
75mm RHA
backblock
0 1.5 1802 16.0 2.58
3203 0 0.7 1355 0.0 2.83
3317 0 0.3 1582 22.5 1.89
3319 0 2.5 1703 29.3 1.88
3318 0 0.8 1814 17.9 2.77
3389 0 0.6 1780 24.0 2.30
3387
30mm SiCB
with coverplate +
75mm RHA
backblock
0 0.4 1464 0.0 3.42
3320 0 1.38 1350 17.00 1.88
3316 0 0.24 1467 14.70 2.49
3314 0 0.65 1574 19.20 2.49
3262 0 1.0 1588 28.4 1.97
3258
30mm SiCB
No coverplate +
75mm RHA
backblock0 2.0 1702 36.1 1.90
3211 0 1.9 1709 16.5 2.82
3376 0 0.9 1590 28.8 1.77
3377 0 0.2 1668 37.76 1.63
3381
40mm SiCB
No coverplate +
75mm RHA
backblock 0 3.6 1788 29.5 2.20
3209 60 1.1 1631 16.6 1.87
3210 60 1.7 1784 25.8 1.85
3244 60 1.5 1689 15.7 2.04
3245
20mm SiCB
No coverplate +
50mm RHA
backblock 60 3.6 1549 0 2.65
3266 60 4.1 1563 20.8 1.71
3253 60 3.2 1672 24.7 1.80
3261 60 5.4 1506 20.1 1.60
3388 60 0.1 1783 35.8 1.67
3386
15mm SiCB
No coverplate +
50mm RHA
backblock60 0.1 1396 0.0 2.43
646 Ceramic Armor Materials by Design
Figure I. 40mm SiC, normal and
oblique impact
0
20
40
60
80
100
120
1000 1200 1400 1600 1800 2000
Impact velocity (m/s)
Re
sid
ual D
OP
into
RH
A(m
m)
RHA reference
CPS/40SiC-B, 0 deg
CPS/40SiC-B, 60 deg
Em=1.3
No-dwell calc.
Figure II. 30mm SiC, normal and
oblique impact
0
20
40
60
80
100
120
1000 1200 1400 1600 1800 2000
Impact velocity (m/s)
Re
sid
ual D
OP
into
RH
A(m
m)
RHA Reference
CPS/30 SiC-B, 0 deg
CPS/30 SiC-B, 60 deg
Em=1.3
No-dwell calc.
Figure III. 30mm & 40mm SiC,
normal impact
0
20
40
60
80
100
120
1000 1200 1400 1600 1800 2000
Impact velocity (m/s)
Re
sid
ual D
OP
into
RH
A(m
m)
RHA Reference
CPS1/30SiCB, 0 deg.
CPS/40SiC-B, 0 deg
EM=1.3 (30mm)
EM=1.3 (40mm)
Figure IV. 30mm & 40mm SiC,
oblique impact
0
20
40
60
80
100
120
1000 1200 1400 1600 1800 2000
Impact velocity (m/s)
Resid
ualD
OP
in
to R
HA
(mm
)
RHA Reference
CPS/30SiC-B, 60 deg
CPS/40SiC-B, 60 deg
EM=1.3 (30mm)
EM=1.3 (40mm)
Figure V. 40mm SiC, normal impact.
Effect of coverplate
0
20
40
60
80
100
120
1000 1200 1400 1600 1800 2000
Impact velocity (m/s)
Re
sid
ual D
OP
into
RH
A(m
m)
RHA Reference
CPS/40SiC-B, 0 deg
NCP/40SiC-B, 0 deg
Em=1.3
Figure VI. 30mm SiC, normal impact.
Effect of coverplate
0
20
40
60
80
100
120
1000 1200 1400 1600 1800 2000
Impact velocity (m/s)
Re
sid
ual D
OP
into
RH
A(m
m)
RHA Reference
CPS1/30 SiC-B, 0 deg
NCP/30 SiC-B, 0 deg
Em=1.3
CPS = Cover plate system
NCP = No cover plate
Ceramic Armor Materials by Design 647
The second reference curve which represents minimal or no rod dwell on
silicon carbide is based on ballistic penetration shots on a second silicon carbide
SiC-100. The shots were performed as an integral part of this current programme
using exactly the same experimental configuration. The targets had no cover
plate. The DOP results from this material yielded a reasonably constant Em of 1.3
across the velocity regime. Previous shock studies on this material [1,2] have
indicated that, on impact, the initial deviatoric strength of this material is
significantly lower than that for SiC B. The quasi-static strengths of the two
materials are almost identical and the density is very similar; 3163 and 3217
kgm-3
for SiC-100 and SiC B respectively.
It is believed that the difference in DOP between SiC100 and SiC B impacted
under identical conditions (compare Em =1.3(30mm) line and the SiC B with no
coverplate data, Figure 6.) is due to the degree of dwell, with SiC-100 exhibiting
little or no dwell. This reference DOP is marked on Figures I-VI as Em =1.3 and
the resulting DOP is calculated for either 30mm thick tiles or 40mm thick tiles.
From Figures I and II it is apparent that the two estimates of zero dwell
penetration estimates do not coincide. It is interesting to note that the ‘no-dwell’
calculation indicates a velocity at which zero penetration terminates is coincident
with the experimental data for both 30mm and 40 mm thick targets. However, the
curve has a very steep slope compared to the Em =1.3 estimate and the
experimental data sets. This may be due to the fact that the penetration rate
measured in reference 4 was measured post-dwell and this could be substantially
different to the penetration rate where no dwell occurred. It was felt that this
estimate is not accurate and the Em=1.3 reference from SiC100 data was adopted
for further comparison.
The residual DOP’s for normal impact and 60 impact angle are compared in
Figure I for a ceramic line-of-sight thickness of 40mm. The reference penetration
of the rod into RHA is shown as a thick solid line. The normal impact data
exhibited a large degree of scatter, with some extremely high Em results (4.4), for
which zero DOP values were measured at high velocities (1770 ms-1
) and some
low Em results (1.6), giving 31 mm DOP at an impact velocity of 1630 ms-1
.
The scatter is much reduced for the oblique targets with Em’s ranging from
1.87 to 2.64. In addition a linear relationship with velocity (correlation coefficient
0.94) is observed for the oblique impact, whilst little correlation is seen for the
normal impact targets. The scatter for the normally impacted targets is broadly
distributed around the oblique impact data suggesting a similar underlying
relationship. When the thickness of the ceramic is reduced to 30 mm both normal
and oblique results follow a very similar trend. Excluding zero penetration values
the Em’s for the normal targets were 2.27±0.44 and for the oblique were 1.69 ±
0.09. At the highest impact velocities the normal 30mm targets deviated from the
oblique producing lower DOP values. It would appear that obliquity does not
offer improved ballistic resistance but tends to reduce the extreme results at both
648 Ceramic Armor Materials by Design
low and high efficiency. Similarly when the target thickness was reduced from 40
mm to 30 mm extreme behaviour was reduced.
It should be noted that for an equivalent LOS thickness at 60 the tile
thickness is halved. One reason for this study was to examine the effect of stress
wave release paths on ballistic performance. Ceramics have high deviatoric
strength under high hydrostatic pressure. When this pressure is released ballistic
performance is reduced. When an oblique target is struck the compressive stress
pulse travels radially from the impact sight but the release will travel back
normally from the back surface of the tile (assuming poor transmission into the
RHA). This means that the release path for the oblique target is half that of the
normal target. The similarity of the normal and oblique results for the thinnest
target (30 mm) would seem to indicate this effect is not dominating the results.
The effect of thickness is plotted in Figures III and IV. Even though there is a
wide distribution in results for the normal impact targets of 40 mm thickness there
is a clear difference between the 30 and 40 mm results. The nominal ‘zero-dwell’
reference lines (Em =1.3, 30 and 40 mm) indicate the difference in DOP that
would result from a purely 10 mm path difference for a constant Em. The 40 mm,
normal impact targets can offer significantly increased performance over the
30mm, exceeding that of purely path difference effects. The possible
improvement is not so great for the oblique targets, even so, there is still an
improvement over and above the path length difference, particularly at lower
velocities. The difference in ballistic performance between 40 and 30 mm normal
impact targets suggests that there are ceramic back surface effects which can
reduce the chances of attaining the extremely high Em’s that SiC B is capable of.
It is difficult to attribute quantitative differences in ballistic performance to
front or back surface effects. The effect of using an impact surface cover plate on
normal incidence targets was investigated for the two thicknesses of ceramic,
Figures V and VI. The cover plate used had a graphite layer adjacent to the
ceramic, used to allow radially spreading rods that were undergoing dwell to
continue easily to spread. For both 40 and 30 mm ceramic targets there appears to
be an increase in ballistic efficiency using the cover plate. For the 40 mm normal
incidence targets there was an apparent improvement in ballistic efficiency at
higher impact velocities. For the 30 mm targets there were improvements in
performance at both high and low velocities. At this stage it is not clear how the
configuration of the cover plate affects the dwell characteristics. In part it may be
due to the reduction of shock impact effects on the surface. It is also possible that
the system promotes dynamic axial confinement of the impacted ceramic surface
by channelling the eroding rod material. Further experiments are continuing with
different configurations.
Ceramic Armor Materials by Design 649
CONCLUSIONS
The factors which promote dwell in silicon carbide are very sensitive to slight
changes in the experimental conditions, resulting in a large degree of statistical
scatter in the determination of ballistic mass efficiency. The effect of obliquity,
thickness and impact surface configuration have been investigated for long rod
impact velocities ranging from 1450 to 1850 ms-1
. Significant benefits in ballistic
performance may be realised by addressing impact surface and ceramic back
surface configurations to maintain reproducibly high ballistic performance.
ACKNOWLEDGEMENT
The work reported in this paper was funded jointly by the UK Government
Corporate Research Programme and by the German Government and was
performed under the auspices of a UK-German collaborative research project.
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650 Ceramic Armor Materials by Design
Ceramic Armor Materials by Design 651
3-D finite element analysis,317
Acoustic impedance, 33Adams, J. W., 629Adams, Marc A., 139Aghajanian, M.K., 527Agrawal, Dinesh, 587Ajayan, Pulickel M., 551Alumina, 63, 83, 91, 103,
185, 233, 269, 441, 463,511, 551, 611, 623, 629
Alumina-zirconia, 91Aluminum nitride, 151Aluminum oxynitride, 573,
587Anderson, Charles E., Jr.,
485Applications, 3
Ballistic performance map(BPM), 139
Bar impact, 225Barker, A.K., 643Bless, Stephan J., 197, 225Bonding, 635Boride, 73Boron carbide, 73, 151, 269Burkett, M.W., 385Burkins, Matthew S., 53, 629
Carbon nanotube, 551Ceramic to metal bonding,
635Chang, Sekyung, 551Chang, Soon Nam, 261, 429Chen, W., 217Chen, Z., 329Chenari, R., 643Cheng, Jiping, 587Chhabildas, Lalit C., 233,
269Cimpoeru, S.J., 361Coated fabric, 541Coating, 541Composite, 73, 185, 551,
611, 623, 629Compressibility, 249Compressive fracture, 197
Compressive layers, 499Computational modeling,
299, 309Cort, G.E., 385Corundum, 463Cost reduction, 451
Damage assessment, 441Damage mechanisms, 557Damage models, 281Dandekar, D.P., 249, 269Danforth, S.C., 473Depth of penetration (DOP),
83, 165, 361Design, 3, 33, 473, 511Doherty, Kevin J., 623, 635Doremus, Robert H., 551Dwell, 113, 173, 309, 557Dynamic fracture, 185Dynamic indentation, 261
Erim, Zeki, 103Ernst, Hans-Jürgen, 23Espinosa, Horacio D., 349
Fabrics, coated, 541Failure mechanism, 103Failure model, 371Fiber, fabric 541Fine grained alumina, 463Finite element analysis, 337,
349Flexible ceramic, 541Forrestal, M.J., 217Fracture mechanics, 185Fragmentation behavior, 103Franks, L. Prokurat, 629Frew, D.J., 217Fused deposition of ceramics,
473Future Combat System, 3Future direction, 421
Gadow, Rainer, 541Galanov, B.A., 73Geometry, 643Gilde, Gary A., 573, 595,
623, 629Glass, plates, 329
Gooch, William A., Jr., 3, 53,113
Grady, Dennis E., 233Grain level analysis, 349Green, William H., 441Grigoriev, O.N., 73Grove, David J., 299, 371
Hbaieb, K., 499High-density ceramic, 45, 53Historical developments, 421Hohler, V., 643Holmquist, Timothy J., 299,
309
Impact surface configuration,643
Impact testing, 113Impact, high-velocity, 23Indentation damage, 429Infra-red windows, 595Interface defeat, 173, 309Isaacs, Jon B., 511Ivanov, S.M., 73
James, Bryn, 33, 165, 643Johnson, Gordon R., 309Joining, 635
Kanel, G.I., 197, 329Kartuzov, V.V., 73Kim, Chang Wook, 261, 429Kim, Do Kyung, 261, 429Kim, Young-Gu, 261Kobayashi, Albert S., 185Kolsky bar technique, 217,
261Konduk, B.A., 103Krell, Andreas, 83, 463
Laminar ceramics, 499Lange, F.F., 499Lanz, W., 63LaSalvia, J.C., 557Layered manufacturing, 473Leavy, Brian, 299Lee, Chul-Seung, 261, 429Lexow, B., 83Lightweight armor, 485
KEYWORD AND AUTHOR INDEX
652 Ceramic Armor Materials by Design
Lischer, David W., 511Lloyd, Isabel K., 623Logan, Kathryn V., 611Long rod penetration, 151,
385Long rod penetrator, 23Lundberg, Patrik, 173
Manufacturing, 91, 451, 473 Marchand, A.H., 385Mashimo, Tsutomu, 233Matthewson, M.J., 473McCuiston, R.C., 473McMeeking, R.M., 499Mears, J., 527Medvedovski, Eugene, 91Membranes, 511Metal-ceramic bonding, 635Meyer, Hubert W., Jr., 299Microcracking diffusion, 329Micro-cracks, 403Micro-mechanisms, 403Microstructure, 349, 557,
611, 629Microwave sintering, 587Modeling, 317, 329, 337,
349, 361, 371, 557Models, comparison of, 299Models, damage, historical
perspective of, 281Molinari, Jean-Francois, 317Morgan, B.N., 527
Nanopowder, alumina, 551Nanotube, carbon, 551Nemat-Nasser, Sia, 403, 511Niesz, D.E., 473Nitride, 73Nondestructive testing, 441Normandia, Michael, 113
Obliquity, 643Orphal, D.L., 151Overview, 3
Palicka, Richard, 53Parker, R., 385Patel, Parimal, J., 573Patterson, Mark C.L., 595Penetration mechanism, 385Penetration model, 337Peron, Pierre-François, 45
Phase equilibrium, 623Pickup, I.M., 643Plane shock wave loading,
249Plastic deformation, 197Polycarbonate, 573Polyurethane, 573Porous silicon nitride, 63Protection areal density
(PAD), 139
Radome, 595Rajendran, A.M., 281, 371Rajendran-Grove model, 371Rao, M.P., 499Rapacki, E.J., 249Razorenov, S.V., 329Reaction bonded silicon car-
bide, 527Reinforcement, 551Reinhart, William D., 233,
269Renström, René, 173Roy, Don W., 595Roy, Rustum, 587Rupert, Nevin L., 441
Sapphire, 233, 573Sarva, Sai, 403, 511Schadler, Linda S., 551Sennett, Michael, 551Shear strength, 249Shear, 557Shen, L., 329Shock compression, 233Shock wave loading, 197Shockey, Donald A., 385Siegel, Richard W., 551Silicon carbide, 63, 73, 151,
269, 309, 441, 527, 635,643
Silicon nitride, 63, 185Singh, J.R., 527Skaggs, S.R., 385Solid freeform fabrication,
473Song, B., 217Spinel, 573, 595Split Hopkinson pressure bar
(SHPB), 217, 269Stassburger, Elmar, 463Stepp, D.M., 421
Strassburger, E., 83Stress propagation, 33Structural ceramics, manu-
facturing, 451Submicron alumina, 83Submicron powders, 463
Target thickness, 643Templeton, Douglas W., 299Test method, 113, 139, 165,
173Tham, R., 643Theory, 139Thermal spray coating, 541Threshold strength, 499Tiles, 33, 103Titanium carbide, 441Titanium diboride, 249, 441,
611, 623, 629Transparent armor, 573, 587,
595Tressler, Richard E., 451Tungsten carbide, 45, 53
Ucisik, A.H., 103Ultra-lightweight armor, 482
von Niessen, Konstantin, 541Vural, Murat, 103
Walker, James D., 337Weber, K., 643Wells, Joseph M., 441Westerling, Lars, 173Wiesner, Volker, 23Wolf, Thomas, 23Wolffe, R.A. 527Woodward, R.L., 361
X-ray computed tomography,441
Zavattieri, Pablo D., 349Zhou, Fenghua, 317Zirconia, partially stabilized
(PSZ), 185